La Metallurgia Italiana

Transcription

N. 4, aprile 2015 - Anno 107
International Journal of the Italian Association for Metallurgy
Organo ufficiale dell’Associazione Italiana di Metallurgia.
Rivista fondata nel 1909
Vi aspettiamo al 22o IFHTSE
Venezia 20-22 Maggio 2015
design: idea-r.it
Poste Italiane spa - Spedizione in abbonamento postale - DL 353/2003 (conv. in L. 27/02/04 n. 46) art. 1., comma1 DCB UD.
Anno 107, n. 4 - Aprile 2015 - Periodico mensile
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European Conference
on Heat Treatment 2015
&
22nd IFHTSE Congress
Heat Treatment and Surface Engineering
From tradition to innovation
VENICE ITALY, 2022 MAY 2015
Organised by
Associazione
Italiana
di Metallurgia
Event no. 395
of the EFC
Sponsored by
Scope
The long-standing co-operation between AIM - Associazione Italiana di Metallurgia - and the IFHTSE - International Federation for Heat Treatment and Surface Engineering - has already led to the joint organisation of
significant events, such as the Congresses held in Florence in 1982 and in 1998 and the Congress on automotive
applications organized in Riva del Garda in 2005.
Now, our combined resources will be focused on the organization of a Conference that will deal with tradition
and innovation in heat treatment and surface engineering. The Conference will join the European Conference
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Programme
Invited lecturer
With about 160 papers in the programme, we expect
ECHT 2015 & 22 IFTHSE in Venice to be a vibrant and
successful event. The 90 oral presentations and over 70
posters will give life to parallel sessions covering:
• Massive heat treatments: microstructural properties
and transformations
• Massive heat treatments: relationship between
process and final properties
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combined processes
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• Quenching media and process
• Massive heat treatments of special components
• Industrial heat treatment equipment and process
innovation
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• Carburizing processes and properties
• Industrial heat treatment equipment and process
innovation
• Applications
• Stainless steels
• Tool steels
• Non ferrous processes and properties
• Coatings properties and characterization
• Coatings processes and applications
• Industrial heat treatment process and simulation
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p.le Rodolfo Morandi 2 · 20121 Milano · Italy
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fax +39 02.7602.0551 · e-mail: [email protected]
conference website: www.aimnet.it/ht2015.htm
www.aimnet.it/ht2015.htm
N 4 apr le 2015 - Anno 107
P ug an
Automa
Organo uffic ale de l Associa ione ta iana di Meta lu gia
R vis a onda a nel 1909
Vi aspettiamo al 22o IFHTSE
Mensile dell’Associazione Italiana di Metallurgia fondata nel 1946
Venezia 20 22 Maggio 2015
7/ 2/04 n 46 a
1 comm 1 DCB UD
g
TAVENGINEERING
Comitato scientifico - Editorial Panel:
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Senk, Piotr R. Scheller, Pierre Soulignac, Jean-Marc Steiler,
Stefano Trasatti, George F. Vander Voort, Maurizio Vedani
i
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7
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02
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os e a a e spa Sp d o e n bb nam n o os a e
nno 07 n 4 p e 20 5 Pe od o mens e
L 353 20 3
onv n L
Direttore Responsabile:
Gianangelo Camona
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FORNO DA VUOTO
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CONS D T V ale Eu opa U i a 29 34073 Gr do GO) Tel 0431 87 070 F x 0431 886507 www co sed t com
nf @c nsed t com
d
N. 4/Aprile 2015
Anno 107 - ISSN 0026-0843
Segreteria di redazione:
Antonella Donzelli
Comitato di redazione:
Federica Bassani, Gianangelo Camona,
Antonella Donzelli, Ottavio Lecis, Carlo Mapelli
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Controllo e caratterizzazione
Comparative analysis on phase quantification methods in
duplex stainless steels weldments
M. Breda, J. Basoni, F. Toldo, C. Bastianello, S. A. Ontiveros
Vidal, I. Calliari ....................................................................3
Storia della metallurgia
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La riproduzione degli articoli e delle illustrazioni è permessa solo citando la
fonte e previa autorizzazione della Direzione della rivista.
Reg. Trib. Milano n. 499 del 18/9/1948.
Sped. in abb. Post. - D.L.353/2003 (conv. L. 27/02/2004 n. 46)
art. 1, comma 1, DCB UD
Consedit sas è iscritta al Roc con il num. 4109
Stampa: Poligrafiche San Marco sas - Cormòns (GO)
Fakes in African art: study of a reliquary figure (MbuluNgulu) from Gabon
C. Soffritti, E. Fabbri, A. Fortini, M. Merlin, G.L. Garagnani ..9
Trattamenti termici
Effetti della diluizione sulla microstruttura e comportamento
ad usura di una lega Fe-C-B-Cr-M
R. Giovanardi, G. Poli, P. Veronesi, G. Parigi, N. Raffaelli ........15
Nanomateriali
Validity of Wulff construction used for
size-dependent melting point of nanoparticles
S. Zhang, L. Zhang, L. Chen .................................................. 25
Acciai
Microstructural characterization and production of high
yield strength rebar
E. Mansutti, G. Luvarà, C. Fabbro, N. Redolfi ...................29
Metalli non ferrosi
On the ageing of a hyper-eutectic Zn-Al alloy
A. Pola, M. Gelfi, G. M. La Vecchia, L. Montesano ...........37
Forgiatura
Implementation of an open-die forging process for large
hollow shafts for wind power plants with respect to an
optimized microstructure
M. Wolfgarten, D. Rosenstock, L. Schaeffer, G. Hirt .........43
Comparative analysis on phase quantification
methods in duplex stainless steels weldments
M. Breda, J. Basoni, F. Toldo, C. Bastianello, S. A. Ontiveros Vidal, I. Calliari
Duplex Stainless Steels (DSS) are biphasic steels of increasing interest, employment as structural materials in
aggressive environments. In these steels, the austenite-to-ferrite phase ratio is maintained at about one – even
if a slightly wider range between 40/60 and 60/40 is in any case accepted – giving the best combination of
mechanical and corrosion-resistance properties. However, DSS must be handled with extreme care, especially
if thermal cycles are involved, owing to the possible formation of dangerous secondary compounds that can
highly worsen their excellent features. In industry, the production of big pipes requires manufacturing welding
operations on steel plates or sheets and the end products must satisfy specific requirements. Therefore,
since DSS properties depend on phase ratio, ferrite quantification at an industrial scale represents a topic of
great interest, which must be as reliable as possible and, at the same time, of fast execution. In the present
paper, different methods currently employed for ferrite estimation in DSS weldments are compared, in order to
understand the limits deriving from each technique.
Keywords: Stainless Steel - Welding - Metallography - Material analysis
INTRODUCTION
In the biphasic austeno-ferritic stainless steels, commonly
named Duplex (DSS), the presence of equal volume fractions of the phases provides an excellent combination of
mechanical properties and corrosion resistance, especially
when compared with conventional stainless steels grades
[1,2]. In Off-Shore engineering, their usage permits the
design of components having smaller thicknesses – and
therefore lighter – without compromising the corrosion resistance and avoiding the employment of expensive anti-
J. Basoni
De Pretto Industrie S.r.l.,
Via Fogazzaro 5, 36015 Schio (VI) - Italy
M. Breda, I. Calliari
Industrial Engineering Department (DII),
University of Padova
Via Gradenigo 6A, 35131 Padova – Italy
F. Toldo, C. Bastianello
Laboratorio Prove Materiali San Marco
Via Lago di Alleghe, 30/32 - 36015 Schio (VI)- Italy
S. A. Ontiveros Vidal
Instituto Tecnológico de Saltillo,
Venustiano Carranza 24000, Tecnologico,
25280 Saltillo, Coahuila de Zaragoza – Mexico
La Metallurgia Italiana - n. 4/2015
corrosion coatings. In DSS, the balanced phase ratio can
be obtained through an appropriate solution-annealing treatment in the temperature range 1050–1100°C followed
by water quenching and, even if the 50/50 ratio is the desired one, phase amounts ranging from 40/60 to 60/40
are in any case accepted.
However, if subjected to improper thermal cycles and
especially in the temperature range 800–950°C, DSS are
sensitive to secondary phases precipitation (intermetallics, carbides, nitrides), which can determine drastic losses
in their advantageous properties [1-8]. Therefore, the solution treatment is usually performed after the forming operations and, besides ensuring the achievement of the desired Duplex microstructure, it permits to re-dissolve any
dangerous precipitate formed during the manufacturing
cycle. However, weldments are not always treatable after
components joining, especially when in-service operations
are performed, and special care must be adopted for welding purposes. In this regard, DSS must be managed as
austenitic grades but using dedicated devices, avoiding
the formation of undesired structures and considering that
the solution-annealing treatment could not be performed
when big parts are joined.
After welding, DSS are required to be free from secondary
phases and the austenite/ferrite volume fractions must be
maintained within the desired forks, in order to guarantee
the expected corrosion resistance and mechanical properties. Therefore, the qualification tests require a systematic
measurement of the ferrite percentage (%FE) in different
parts of the joint – base material (BM), heat affected zone
(HAZ) and fusion zone (FZ). The international reference
3
Memorie
C
Si
Mn
Cr
Ni
Mo
Cu
P
S
N
forged
0.027
0.22
1.12
22.2
5.9
3.3
0.16
0.023
0.002
0.16
HIP
0.021
0.71
1.13
22.4
5.2
3.1
0.17
0.018
0.003
0.17
Table 1 – Chemical composition of the base materials [wt.%].
standards assess how to evaluate %FE, but customers
often necessitate dedicated procedures by adapting the
standardized approach to their specific requirements. The
most widely accepted international standard for phases
quantification is Optical Microscopy (OM) after polishing
and chemical etching, by following the provisions of ASTM
E562-11 standards (manual point count), which provides
information on selecting type of pattern and number of
fields to be analyzed that ensure a well-defined relative
accuracy. On the other hand, points counting through automatic image analysis is otherwise possible, in accordance with the provisions of ASTM E1245-03 standards (automatic image analysis). In both cases, the use of Scanning
Electron Microscope (SEM) instead of OM is permitted.
Phase quantification by image analysis on OM and SEM
micrographs is often replaced by other simple – and faster
– field-methods, among which the use of ferritoscope is
the most popular. This technique is based on measuring
the magnetic field generated by an induced-currents probe, by placing it in contact with the metal surface; since
ferrite is a magnetic phase while austenite is amagnetic,
this device provides %FE by measuring the magnetic response of the material. However, this method is very sensitive to the finishing of the contact surface and cannot be
applied near edges or corners, owing to the distortion of
the magnetic field. Moreover, the use of ferritoscope is limited to the investigation of wide areas of welds and base
material, whereas HAZ is not readily controllable, due to
its small size. In this method, results are given in units called Ferrite Number (FN) and are automatically converted
to %FE through an internal correlation.
Another method for %FE determination is based on calculations from phase diagrams, according to the AWS 5.4
assessments, by knowing the Cr equivalent (Creq) and Ni
equivalent (Nieq) values of both base and filler materials
derived from chemical analyses. A disadvantage of this
method lies on the low accuracy level of the obtained values, since the lines on diagrams are drawn only for some
reference values and all points in between are evaluated
graphically, implying many difficulties and a dependence
on the adopted interpolation method. Moreover, if on one
side BM and FZ compositions can be easily determined,
the chemical analysis on HAZ cannot be executed, owing
again to its small size.
In the present paper, a comparison of the previously described method for %FE evaluation on SAF 2205 DSS welded joints is reported, in terms of relative reliability and
associated accuracy.
4
EXPERIMENTAL
The comparative study was carried out on two full-penetration qualification beads (Bead-1 and Bead-2), obtained
by joining two types of SAF 2205 DSS (UNS S31803) base
materials and adopting the same welding procedure. The
base materials (compositions in Table 1) were of different
manufacture, produced by traditional forging (forged material) and Hot Isostatic Pressure (HIP material), whereas
the welding procedure involved the Gas Tungsten Arc Welding (GTAW) process for the first passes (about 6 mm of
deposited material), subsequently filled using the Shielded
Metal Arc Welding (SMAW) technique.
The qualification beads were prepared for metallographic
investigation on a Leica DMRE OM by mechanical polishing
and electrolytic etching at 5V using a solution composed
of 30% NaOH in deionized water. In each sample, different zones of the weldments were distinguished: BM (at
about half thickness), upper HAZ (cover layers), lower HAZ
(root layers), upper FZ and lower FZ. For the comparative analysis, %FE was estimated taking into account four
different methods: ASTM E562 manual counting points
method using fixed parameters (magnification, grid and
number of fields), ASTM E562 manual counting points method using variable parameters, automatic image analysis
(ASTM E562 and ASTM E1245) and manual image analysis
(variable magnification). For the first three OM methods,
a 500x magnification for all the investigated areas was
maintained.
The same zones of the weldments were observed using a
Leica Cambridge Stereoscan 440 SEM operating in backscattered-electron mode (BSE) at 29 kV; as it is known,
the SEM-BSE observation allows distinguishing the microstructural constituents according to their average atomic
number and, therefore, ferrite (the lighter phase) appears
darker than austenite (the heavier phase). In this case, the
samples were slightly etched, in order to only better define
phase boundaries, while the magnification was set as variable, according to the observed microstructure. For ferrite quantification, the micrographs were edited using an
image-analysis software by applying proper filters to improve phase contrasts and minimize grayscale threshold
errors (the filtering procedure was set-up for the specific
case).
For the aim of the present work, %FE was also measured
using a portable Fischer MP30 Ferritscope in the previously defined areas, and calculated using phase diagrams
(Schaeffler-Delong, ESPY and WRC-1992), after the evaluation of the local chemical composition and considering
the Creq and Nieq values.
Fig. 1 – OM images of
etched microstructures
in the investigated DSS
weldments (Bead-1):
forged BM-HAZ (left), FZ
(middle), HIP HAZ-BM
(right) .
RESULTS AND DISCUSSION
In both the qualification beads, the microstructures were
similar; the BMs were free from intermetallics and the welded zones exhibited the classical dendritic morphology
achieved from melt solidification (Figs. 1 and 2). As expected, the welding processes (GTAW and SMAW) determined
some differences in the final FZ microstructure: SMAW,
owing to its lower heat input, caused the formation of a
finer dendritic structure in the filling passes, but an increased micro-porosity level respect to GTAW was obtained.
The results of the investigation are listed in Table 2, where
the values relative to Bead-1 are reported over those concerning Bead-2. Starting from the OM quantification, it is
possible to note that increasing the number of grid points,
the accuracy of the evaluation also increased, as the standard deviation was reduced; however, a major number of
points lead to an increase in quantification time spent by
the operator, which is not suitable at an industrial scale.
The variable magnification method did not provide improvements on phase quantification, and the results were similar to those obtained using the manual method with a
low number of grid points. On the contrary, the automatic
image analysis software provided the more accurate results, even using a smaller number of fields, because the
system is able to automatically delete the “problems” concerning the choice of the suitable grating; however, in this
case, the evaluation procedure must be properly set-up.
As can be seen from the table, the middle part of FZ was
not always taken into account, since it is a transient region
and requires particular assessments. Concerning the OM
automatic method, the micrographs can be easily edited
using image-analysis, because the employed reagent darkens ferrite and leaves austenite unaffected; therefore, the
images appear in a grayscale pointing toward a black-andwhite image, thus having a net phase contrast, allowing
for a simple determination of ferrite volume fraction and
easing phases discrimination.
The analysis of SEM images showed substantial differences respect to OM manual quantification, especially in the
base materials, and ferrite was mainly underestimated. In
this regard, SEM introduces a further source of uncertainty, since the signal from the BSE detector is processed
by assigning different levels of grey to the detected energies. Therefore, the micrographs are no more like blackand-white images, but an extended greyscale image is
created, and the threshold defining the boundary between
the phases cannot be univocally defined at all (Fig. 2). In
this case, the operator plays a key role in the accuracy
of the estimated phase amount, because the assignation
of the threshold became strongly subjective. In addition,
austenite and ferrite compositions in FZ are conditioned
by the welding processes, since temperature and cooling
rate affect elements partitioning. Therefore, changes in
average atomic number can occur, leading to difference
5
Memorie
Method
%FE (st.dev.)
Forged BM
Upper HAZ
Lower HAZ
Upper FZ
Middle FZ
Lower FZ
HIP BM
manual OM
16 pt. grid
(500x)
52 (12)
51 (7)
60 (13)
58 (8)
55 (9)
57 (6)
37 (9)
41 (10)
-
30 (6)
30 (9)
-
manual OM
96 pt. grid
(500x)
52 (5)
50 (6)
61 (3)
62 (4)
57 (5)
58 (5)
32 (3)
37 (4)
-
30 (3)
31 (3)
-
automatic
OM
(500x)
54 (1)
52 (1)
59 (1)
59 (2)
60 (1)
56 (2)
34 (2)
38 (1)
55 (4)
-
32 (2)
32 (1)
33 (9)
-
manual OM
(var. magn.)
50 (7)
56 (10)
-
-
36 (3)
43 (5)
46 (3)
37(6)
29 (4)
34 (8)
56 (2)
57 (2)
SEM
(var. magn.)
44 (6)
47 (7)
58 (7)
64 (8)
59 (8)
65 (9)
38 (4)
43 (6)
40 (4)
36 (9)
26 (5)
32 (7)
49 (3)
52 (2)
Ferritoscope
53 (1)
58 (1)
-
-
37 (1)
40 (2)
-
32 (1)
26 (1)
-
*Schaeffler
49
-
-
38
34
-
-
49
ESPY
82
-
-
28
25
-
-
89
*WRC-1992
84
-
-
56
51
-
-
98
* Value in FN
Table 2 – Comparison of ferrite quantification methods (Bead-1 over Bead-2).
Fig. 2 – SEM images of etched microstructures (Bead-1): forged HAZ (left), upper FZ (middle), lower FZ (right).
in the observed microstructure. In the present study, the
quantification was slightly facilitated by the preliminary
electrochemical etching, but the introduction of a large
amount grayscales unavoidably leaded to an increase in
estimation errors.
Ferritoscope provided values that sometimes coincided
and sometimes differed from those obtained through
image methods; this was mainly due to the extent of the
analysed volume (about 10 mm3), greater than that involved using other techniques, and also to the amplification
of the uncertainty in the conversion from FN to %FE. Finally, the values calculated from phase diagrams showed the
greatest variability, which is related to the chosen method.
6
Except for the Schaeffler-Delong diagram, the results were
highly misaligned to that of OM, SEM and ferritoscope;
this is intrinsically due to such diagrams, which have been
developed to perform calculations on areas having a high
concentrations of alloying elements, more similar to that
of filler materials rather than the welded ones.
CONCLUSIONS
In the present paper, different methods currently employed for ferrite quantification in DSS weldments were compared, in order to understand the limits related to each
technique. Analyses involving OM, SEM, ferritoscope and
phase diagrams were considered, and the related ferrite
estimations were presented.
From the results, it was revealed that OM is preferable for
ferrite quantification if compared to SEM, both in base material and in the welded zone. In manual counting, several
grid points must be adopted for a proper phase estimation
and the uncertainty can be strongly reduced, but the increased counting time makes this method not readily applicable when fast estimation are required. On the other
hand, the automatic method provided the best results,
but the procedure must be properly set-up; in this case,
OM images are nearly black-and-white and, therefore, the
subsequent analysis can be easily performed, since the threshold values are better defined and the “problems” concerning the choice of the suitable grating are automatically
deleted by the program. On the contrary, SEM micrographs are grayscale images and phases quantification highly
suffers from the assigned thresholds by the operator to
phase boundaries. Moreover, since phases compositions
are conditioned by the previous thermo-mechanical processes, difference in the observed microstructure may occur by adopting the SEM-BSE observation method, owing
to changes in the phase average atomic number.
Finally, the investigation confirmed the low accuracy of
fields-methods such as those deriving from phase diagrams, whereas ferritoscope can be considered as an
intermediate-accuracy technique, even if it requires large
volumes to be measured. These methods, although simple
and fast, cannot always assure a reliable ferrite quantification, owing to uncertainties intrinsic to the methods itself.
REFERENCES
[1] J. O. Nilsson. Super Duplex Stainless Steels. Mater Sci
Tech 8 (1992), p. 685.
[2] R. N. Gunn. Duplex Stainless Steels: Microstructure,
Properties and Applications. Abington Publishing,
Cambridge, England (1997).
[3] J.O. Nilsson, A. Wilson. Influence of isothermal phase
transformations on toughness and pitting corrosion of
super duplex stainless steel SAF 2507. Mater Sci Tech
9 (1993), p. 545.
[4] I. Calliari, M. Zanesco, E. Ramous. Influence of isothermal aging on secondary phases precipitation and toughness of a duplex stainless steel SAF 2205. J Mater
Sci 41 (2006), p. 7643.
[5] I. Calliari, G. Straffelini and E. Ramous. Investigation of
secondary phase effect on 2205 DSS fracture toughness. Mater Sci Tech 26 (2010), p. 81.
[6] I. Calliari, M. Pellizzari and E. Ramous. Precipitation of
secondary phases in super duplex stainless steel Zeron100 isothermally aged. Mater Sci Tech 27 (2011),
p. 928.
[7] I. Calliari, M.Breda, E. Ramous, K. Brunelli, M.Pizzo, C.
Menapace. Impact Toughness of an Isothermally Treated Zeron®100 SDSS. J Mater Eng Perform 21 (2012),
p. 2117.
[8] M. Pohl, O. Storz, Thomas Glogowski. Effect of intermetallic precipitations on the properties of duplex
stainless steel. Mater Charact 58 (2007), p. 65.
7
11th EuropEan
ELECtrIC StEELMaKInG
ConFErEnCE & Expo
Venice (Italy) • 25-27 May 2016
Organised by
AssociAzione
itAliAnA
di MetAllurgiA
Patronised by
www.aimnet.it/eec2016.htm
After the successful events in Florence (1986) and Venice (2002), AIM hosts the Electric Steelmaking Conference & Expo 2016.
The total world steel production coming from electric arc furnaces is today up to 35% and there is a ceaseless technical development
of EFC process and of its auxiliary equipment. In addition, the combination of such technology with new plants based on direct
reduction of iron ores makes electric furnace an interesting solution not only for the production route based on steel scrap recycling
but also for a flexible usage of mineral resources.
The 11th European Electric Steelmaking Conference & Expo aims at providing a forum for information transfer of the latest
techniques and applications in electric furnace steelmaking and at bringing together steel producers, research and academics
active in steelmaking industries, energy production and plant supply concerns.
◗◗◗ Technical Focus
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Island of San Giorgio Maggiore in Venice.
◗◗◗ Exhibition & Sponsorship
As an integral element, EEC 2016 will feature an Exhibition that will enable
excellent exposure for company products, technologies, innovative
solutions or services. At this opportunity the Organizers will set an area
focal point of the Conference, so as to guarantee a perfectly targeted
potential customer’s environment. Companies will also be able to reinforce
their participation and enhance their corporate identification by taking
advantage of benefits offered to them as Sponsors of the Conference.
Companies interested in exhibition & sponsorship may contact the
Organising Secretariat (e-mail: [email protected] / fax: +39 0276020551).
◗◗ Important dates
Deadline for submission of abstracts
Information on Acceptance
Opening of the online registration
Deadline for Full Paper Submission
May 31, 2015
October 31, 2015
October 31, 2015
January 31, 2016
AIM is looking forward to welcoming you in the unique city of Venice!
EEC 2016 Organizing Secretariat
AIM - Associazione Italiana di Metallurgia ⋅ P.le R. Morandi, 2 20121 Milan - Italy ⋅ Tel. +39 02 76021132 ⋅ Fax +39 02 76020551
E-mail: [email protected] ⋅ Website: www.aimnet.it/eec2016.htm
Storia della Metallurgia
Fakes in African art: study of a reliquary
figure (Mbulu-Ngulu) from gabon
C. Soffritti, E. Fabbri, A. Fortini, M. Merlin, g.L. garagnani
The aim of the present work is the chemical and microstructural characterisation of a reliquary figure, stylistically
consistent with the art of the Kota population, which lived in the eastern part of Gabon (Africa). The artefact was
subjected to preliminary observation by stereomicroscopy, and then Optical Microscopy (OM) and Scanning
Electron Microscopy (SEM) analyses are carried out on a fragment and on surface compounds. Lastly, AMS
radiocarbon dating of the wooden support allowed further information about the production period to be obtained.
The results show that the artefact was produced by a Cu-Zn alloy and contains non-metallic impurities made up of
S and Se. The greenish and whitish surface compounds, which are mainly collected near the nails and in proximity
to the overlaid sheets, are probably only partly related to natural corrosive processes. Finally, radiocarbon dating
established that the wooden support certainly dates after 1950.
Keywords: Copper and alloys – Material characterisation – Metallography – Electron microscopy – Metallurgy
INTRODUCTION
The proliferation of fakes in African arts has grown enormously in recent years, with a particular explosion since
the 1950s, due to an increase in demand by collectors,
which created new fields of activities for African foundries. In fact, in the 1980s the quantity of antiquities on sale
increased further and today many replicas of tourist souvenirs and fanciful copies of traditional forms enrich the
art market.
The official definition of authenticity for African artefacts
consists of two inseparable conditions: any object created
for a traditional purpose and by a traditional artist may be
considered authentic [1].
It is rather difficult to determine if an African artefact is
original or a copy because literature is characterised by
incomplete information about the African arts and the production of artefacts by artists [2]. The studies of African
artefacts are somewhat incomplete since there is no cor-
C. Soffritti
TekneHub, Department of Architecture, University of
Ferrara, Via Quartieri 8, 44121 Ferrara, Italy
C. Soffritti, E. Fabbri, A. Fortini, M. Merlin,
g.L. garagnani
Department of Engineering - “A. Daccò” Corrosion and
Metallurgy Study Centre, University of Ferrara,
Via Saragat 1, 44122 Ferrara, Italy
relation between the style used in these works, the materials used to produce them and the geological context of
the extraction zone.
Archaeometric analyses are essential to determine the
state of conservation of the objects as well as to evaluate
the production period in order to establish the authenticity
of the artefacts.
The aim of the present work is the characterisation of a
sculpture, stylistically consistent with the art of the Kota
population, which lived in the eastern part of Gabon (Africa). This community is known for the realisation of metallic reliquary figures, which were set on wooden supports
and called Mbulu-Ngulu or Bwéte. It should be noted that
the first samples of these sculptures arrived in France and
Germany during the last quarter of the 19th century. Reliquary artefacts should, however, be more ancient given
that the local copper mines had already been exploited to
obtain metal for the coating of artefacts [3].
Unfortunately, the majority of researchers have relied solely on stylistic analyses of the ornaments, which decorate
the surface of the objects. More extensive investigations
on the chemical composition of the alloy of the artefacts,
in conjunction with a systematic characterisation of original African metallic objects would allow the evaluation of
the provenance and the dating of Kota funeral art [4].
The present paper focuses on the investigation of the
symbolic representation of a human abstract figure whose
head is bigger than the rest of the body. These abstract
figures were used to protect and demarcate the bones
of family ancestors, which were preserved in containers
made of bark.
The artefact consists of a carved piece of wood (42 cm in
height, 23 cm wide and 2 cm thick) covered on one side
with metal sheets, which were fixed onto the support with
9
Memorie
Fig. 1 – Macroscopic images
of the sculpture: front (a),
back (b) and side view(c).
Fig. 1 – Immagini fotografiche
del manufatto: parte anteriore
(a), parte posteriore (b) e
visione laterale (c).
small metallic nails. These metal sheets are very thin in
order to fix almost perfectly to the carved wood. The following morphological elements of the sculpture are detectable in Fig 1a-c:
• The oval face has stylised eyes and nose but the mouth
is not depicted. Two metallic plates are nailed onto the
surface to represent a cross. In agreement with E.W.
Herbert [5] this element has been found starting from
the end of the 15th century as a result of the Congolese
population’s conversion to Catholicism;
• Two lateral parts at ear-level which are often considered the representation of a hat;
• Two cylindrical pendants placed on the base of the lateral parts, which are the abstract representation of
traditional male and female hairstyles;
• One half-moon shaped sheet is located above the
oval face and harmoniously integrated with the lateral
parts;
• One rhomboid element that symbolises the body and
the legs, placed on a rectangular wooden base;
• Another rhomboid element, on the back of the sculpture, stretched along the vertical axis and with a protruding “vein”.
The goal of the present work is the chemical and microstructural characterisation of the sculpture as well as of
the products located on its surface. The analyses were
carried out by stereomicroscopy, Optical Microscopy (OM)
and Scanning Electron Microscopy (SEM) coupled with
Energy Dispersion Spectroscopy (EDS). Finally, a wooden
fragment of the support was analysed using Accelerator
Mass Spectrometry (AMS), which enabled radiocarbon
(14C) dating of the artefact.
Materials and methods
The sculpture was observed by stereomicroscopy, equipped with a Moticam 2500 – 5 Mp camera, in order to obtain information on the manufacturing technique and to
check the state of conservation. The investigations have
revealed the presence of some compounds, which are
10
mainly concentrated near the nails and in proximity to the
overlaid sheets.
Thereafter, the evaluation of the alloy and the composition
of the different colour surface compounds was carried out
using a ZEISS EVO MA 15 Scanning Electron Microscope (SEM), coupled with Energy Dispersion Spectroscopy
(EDS).
Moreover, a metal fragment of a few millimetres was taken
from an unobtrusive area of the sheet. The sample was
mounted in conductive resin, polished and submitted to
conventional metallographic observation using LEICA MEF4M Optical Microscopy (OM).
Lastly, a sample of a few grams was collected from the
base of the wooden support and subsequently was dated
using Accelerator Mass Spectrometry (AMS) at Centro di
Datazione e Diagnostica (CEDAD) – University of Salento.
Results and discussion
Macroscopic investigations
Preliminary macroscopic investigations have yielded a great deal of information about the manufacturing technique
as well as the nature of the products located on the surface.
In agreement with some of the literature and private communications expressed in the last few years [4], the parts
of the face that are not overlaid by the two metallic sheets
(positioned in a cross) consist of a single plate. These are
decorated with “lamellage”, a technique characterised by
various equidistant streaks, which are placed in a slanting
or horizontal pattern. These strips are also depicted on
the half-moon sheet and on the rhomboid element that
symbolises the body and the legs. In the latter two cases,
a multitude of pitting embossed using a punch also decorates the surface of the plate [4].
Fig. 2 shows representative images of the rare compounds
using stereomicroscopy, which were mainly collected near
the nails and in proximity to the sheets. The colouring of
these products is clearly green or whitish.
In particular, in Fig. 3a a shrinkage cavity of remarkable
dimensions, formed during the alloy solidification, is visible. Fig. 3b also shows the microstructure of the alloy
after chemical etching by FeCl3/HCl. The presence of
both non-homogeneous grain size and thermal twin bands
would suggest that the artefact was obtained by alternate
hammering and annealing steps. It should be noted that
the variable grain size is probably due to a heterogeneous
plastic deformation induced by manual hammering.
Fig. 2 – Macroscopic images of the different colour
compounds that appear on the surface.
Fig. 2 – Macrografie rappresentative dei composti che si
presentano con cromie differenti.
Chemical analysis
Fig. 4a shows a SEM image of the alloy together with the
corresponding EDS spectrum. SEM-EDS analysis highlights
that the artefact was produced by a Cu-Zn alloy, without
the addition of alloying elements, i.e. Pb. No impurities
(i.e. As, Fe, Sb), which are very common in the ancient
alloys, were detected. It should be noted that, comparable amounts of Cu and Zn (Fig. 4b) could also be found in
modern brasses such as the commercial “Yellow Brass”
which contains 65 wt.% of Cu and 35 wt.% of Zn [7].
Fig. 5a shows a SEM image of rare microscopic inclusions
that are visible in the alloy. In particular, SEM-EDS analy-
Fig. 3 – Optical images of: (a) a detail of a shrinkage cavity on the polished surface; (b) the microstructure of
grains with the presence of thermal twin bands on the chemically etched surface.
Fig. 3 - Micrografie OM del manufatto: cavità da ritiro (a), in assenza di attacco metallografico; microstruttura a grani e
geminati (b); in presenza di attacco con reattivo a base di cloruro ferrico.
Radiocarbon measurement
The AMS radiocarbon dating established that the wooden support is certainly dated after 1950. It is well known
that the production of these sculptures ended around the
1930s because of the great number of Catholic missions,
which imposed a new social organisation based on Western households [4]. Finally, it should be pointed out that
the integrity as well as the total absence of signs due to a
wooden support substitution is clearly evident.
Microstructural analysis
After metallographic preparation, the microstructure of
the metal fragment taken from the sheet was highlighted.
sis allows the verification of the presence of non-metallic
impurities enriched with S and Se. To our knowledge, only
one reference reports some South African (Lowveld) metallic artefacts [8], approximately dated from 1000 A.D. to
1980 A.D., which were characterised by many copper-iron
sulphide inclusions containing up to 3% Se by weight, residual from incomplete ore reduction.
Over the centuries, the Kota reliquary figures went into
stylistic decline and they were characterised by more abstract and grotesque meanings. Moreover, the demand
for these artefacts from Western collectors has grown
enormously in recent years, causing the proliferation of
sculptures without any “funerary” meaning for the purpo-
11
Memorie
Fig. 4 – SEM backscattered electron image of the
alloy indicated by pink square (Spectrum 1), together
with the corresponding
EDS spectrum; (b) average
composition of the area in
Fig. 4a (measured by EDS).
The contents of Cu and Zn
are highlighted in Fig. 4b
Fig. 4 – (a) immagine al
SEM della matrice metallica
con indicazione (Spectrum
1) della zona analizzata e
corrispondente spettro EDS;
(b)
dati
semi-quantitativi
relativi allo spettro in Fig. 4a.
In Tabella vengono evidenziati i contenuti di Cu e Zn.
Element
Weight % Atomic %
CK
2.44
11.59
OK
0.57
2.03
Cu K
62.65
56.35
Zn K
34.35
30.03
Totals
100.00
Element
Weight % Atomic %
CK
12.17
40.42
SK
9.38
11.66
Cu K
36.53
22.92
Zn K
36.60
22.32
Se K
5.32
2.69
Totals
100.00
Fig. 5 – SEM backscattered electron image of an inclusion, indicated
by a black arrow (Spectrum 2), together with the corresponding
EDS spectrum; (b) average
composition of the analysed
point in Fig. 5a (measured by
EDS). The contents of S and
Se are highlighted in Fig. 5b.
Fig. 5 – (a) immagine al SEM
di una delle inclusioni visibili
all’interno della matrice metallica
con indicazione (Spectrum
2) della zona analizzata e
corrispondente spettro EDS;
(b) dati semi-quantitativi relativi
allo spettro in Fig. 5a. In Tabella
vengono evidenziati i contenuti
di S e Se.
Fig. 7 – Representative SEM image of the morphology of surface whitish compounds, together with
corresponding EDS spectrum.
Fig. 7 – Micrografia SEM dei composti di colore biancastro e relativo spettro EDS.
12
se of enriching the flourishing art market. Starting from
the first decade of the 20th century, the practice of recasting damaged copper and brass to recover the precious
metal was very common. In particular, E. Andersson [9]
highlighted that many “Mbulu-Ngulu” were obtained by
recasting ancient alloy which was later mounted in more
recent wooden supports (second half of the 20th century).
In this regard, because of the perfect realisation of the
artefact without the addition of alloying elements or common impurities, the rare Se inclusions and the dating of
the wooden support (see § Radiocarbon measurement), it
is possible to further suppose that the sculpture analysed
in this paper was realised by the methods described by
Andersson in [9].
The SEM image of the greenish surface compounds together with the corresponding EDS spectrum is reported in
Fig. 6. Because of the small amount of products on the
surface, it was not possible to take samples and to carry out specific analyses like XRD or Raman spectroscopy.
First of all, the morphologies in Fig. 6a and 6b are very different. It should be noted that the needle-like or lamellar
structure shown in Fig. 6a is frequently observed in copper carbonate compounds. This evidence is supported by
SEM-EDS analysis. On the contrary, the same technique
would suggest that the compounds of Fig. 6b are probably
zinc oxychloride.
Fig. 7 is a representative SEM image of whitish compounds, which are mainly collected in proximity to the overlaid
sheets. The EDS spectrum emphasises high concentrations of Cl and Pb. In particular, the latter element is totally
absent in the alloy and it is possible that it is not produced
by natural corrosive processes.
BIBLIOGRAPHY
[1] J. CORNET, Afr. Arts, Vol. 9 (1975), no. 1, pp. 52-55.
[2] E. BASSANI, Rivista trimestrale di studi e documentazione dell’Istituto italiano per l’Africa e l’Oriente (1980),
no. 1 pp. 85-95.
[3] E. BASSANI, Arte Africana, Ed. Skira, Milano (2012).
[4] G. DELORME, Arts d’Afrique Noire, (2002), no. 122, pp.
1-14.
[5] E.W. HERBERT, Red gold of Africa - Copper in precolonial history and culture, University of Wisconsin Press,
Madison (1984).
[6] E. ANDERSSON, Contribution à l’ethnographie des
Kuta, Ed. Almqvist & Wiksell, Stockholm (1953).
[7] ASM Specialty Handbook - Copper and Copper Alloys,
ASM International, Materials Park, Ohio, (2001).
[8] D. MILLER, D. KILLICK, N.J. VAN DER MERWE, J. Field
Archaeol., Vol. 28 (2001), pp. 401-417.
[9] E. ANDERSSON, Contribution to the ethnography of
Kuta III, Occasional Papers XV, S. Lagercrantz and A.
Loôv (1991).
Concluding remarks
The present work has proved the usefulness of an interdisciplinary approach to clarify some general aspects about
the manufacturing process and the state of conservation
of metal artefacts.
Macroscopic examinations have highlighted a good state
of conservation of the sculpture and a manufacturing process consistent with the reliquary Kota art.
Observations by Optical Microscopy (OM) have established that the sculpture was obtained by a casting and was
subsequently subjected to alternate hammering and annealing stages.
SEM-EDS analysis has highlighted that the artefact was
produced by a Cu-Zn alloy, with an amount of the latter
elements comparable to those that could can be found in
modern brasses (i.e. “Yellow Brasses”). The absence of
alloying elements and the presence of rare Se inclusions
bear witness to an advanced manufacturing process and
this suggests that the artefact was obtained by a rather
recent recast. Finally, the chemical analyses of greenish
and whitish surface compounds lead to the assumption
that they are only in part related to natural corrosive processes.
13
STEELSIM
2015
6th International Conference
MODELLING and SIMULATION of METALLURGICAL PROCESSES in STEELMAKING
BARDOLINO, ITALY
23-25 SEPTEMBER 2015
www.aimnet.it/steelsim2015.htm
Modelling and Simulation of metallurgical processes cover an important role in optimizing technological processes, decreasing
production costs, increasing steel quality and defining the correct design of metallurgical processes in order to improve their
sustainability even from the environmental point of view.
The fundamentals of metallurgical processes can be investigated through physical and numerical modelling following several
numerical approaches.
Traditional and new mathematical techniques applied by modern simulation facilities allow to achieve results that are useful
to understand physical interaction and to design a profitable metallurgical process.
The simulations technique can be applied to the different steps of the metallurgical production route: production and refining
of liquid metals, solidification, plastic deformation, thermo-mechanical processes, thermal treatment, verification of structural
reliability etc.
●●● CONFERENCE TOPICS
State of art and developments in modeling and simulation
in steelmaking:
• Ironmaking
• Primary metallurgy (aluminium alloys, copper alloys
titanium alloys etc.)
• Secondary steelmaking
• Refining of metal alloys
• Thermodynamic and kinetic simulation of the metallurgical
systems
• Casting and solidification
• Electrochemical processes
• Metalforming processes and thermo-mechanical treatment
• Heat treatments
• Fracture mechanics and safety criteria
• Fatigue mechanics
• Safety criteria
• Reduction of environmental impact
●●● EXHIBITION & SPONSORSHIP OPPORTUNITIES
SteelSim 2015 will feature an Exhibition that will enable
excellent exposure for company products, technologies,
innovative solutions or services.
Companies will also be able to become Sponsors of the
Conference.
Companies interested in taking part in the Exhibition or in
sponsoring the event may contact the Organising
Secretariat (e-mail: [email protected] / fax: +39 0276020551).
●●● VENUE
The Conference will be staged at the Congress Center
of Aqualux Hotel Spa Suite & Terme, in Bardolino (VR), Italy
Via Europa Unita, 24/b 37011
Organised by
ASSOCIAZIONE ITALIANA DI METALLURGIA
ORGANIZING SECRETARIAT
AIM - ASSOCIAZIONE ITALIANA DI METALLURGIA
P.le R. Morandi, 2 · 20121 Milano · Italy · tel. +39 02 76021132 · fax. +39 02 76020551
e-mail: [email protected]
www.aimnet.it
Trattamenti termici
Trattamenti termochimici di nitrurazione
e post-ossidazione su acciai 17-4PH:
ottimizzazione dei parametri di processo per
massimizzare la resistenza a corrosione
R. giovanardi, g. Poli, P. Veronesi, g. Parigi, N. Raffaelli
L’acciaio inossidabile 17-4PH viene solitamente trattato termicamente per incrementarne le proprietà
meccaniche. Per migliorare ulteriormente la resistenza ad usura di tale acciaio è possibile sottoporlo a trattamenti
termochimici, quali ad esempio la nitrurazione. Trattandosi di un acciaio inossidabile, in grado di presentare allo
stato di fornitura una notevole resistenza a corrosione conferita dall’elevato contenuto di cromo presente in lega,
viene spontaneo chiedersi se e quanto i trattamenti termici e termochimici possano influenzare questa proprietà.
Lo scopo del presente lavoro è quello di valutare come l’applicazione di trattamenti termici e termochimici,
solitamente impiegati per incrementare proprietà meccaniche e anti-usura, influiscano sulla resistenza a
corrosione dell’acciaio e di intervenire sulle variabili di processo degli stessi (oppure mediante trattamenti
successivi quali la post-ossidazione) al fine di individuare le condizioni di trattamento ottimali per preservare
una discreta resistenza a corrosione. A tale scopo sono state eseguite prove di corrosione accelerata, mediante
acquisizione di curve di polarizzazione in cella elettrochimica, su provini sottoposti a diverse combinazioni di
trattamenti termici e termochimici (invecchiamento H1025, nitrurazione, post-ossidazione) eseguiti in diverse
combinazioni di tempi e temperature. Oltre alla caratterizzazione elettrochimica i provini sono stati sottoposti a
prove di microdurezza HV superficiale ed in sezione, per valutare l’effettiva efficacia dei trattamenti applicati in
termini di proprietà meccaniche ed antiusura. Nonostante i migliori risultati in termini di incremento della durezza
superficiale e di profondità di indurimento siano stati raggiunti con trattamenti che compromettono notevolmente
la resistenza a corrosione dell’acciaio, il lavoro svolto ha permesso di individuare ed ottimizzare sequenze di
trattamenti che permettono di preservare quasi completamente la resistenza a corrosione dell’acciaio, pur
incrementando la durezza superficiale fino a valori di oltre 850HV.
Parole chiave: Acciaio inossidabile - Corrosione - Trattamenti termici - Caratterizzazione materiali
INTRODUZIONE
Gli acciai inossidabili PH (precipitation hardening) trovano
impiego in una varietà di applicazioni, quali raccordi aerei
[1], ingranaggi e fasteners [2, 3], componenti per reattori
nucleari [4-7], tuttavia il loro utilizzo è fortemente limitato
dalla loro relativamente bassa durezza e soprattutto dalR. giovanardi, g. Poli, P. Veronesi
Università di Modena e Reggio Emilia,
Dipartimento di Ingegneria ‘Enzo Ferrari’,
Via Vignolese 905, 41125 Modena
g. Parigi, N. Raffaelli
STAV srl
Via della Lora 18/I-N,
50031 Barberino del Mugello (FI)
le scarse proprietà tribologiche. Per questo motivo risulta
estremamente interessante la possibilità di applicare trattamenti superficiali o rivestimenti che possano incrementare le proprietà antiusura di tali acciai. In questo lavoro
verranno sperimentati trattamenti termochimici di nitrurazione e post-ossidazione sull’acciaio 17-4PH, con lo scopo
di incrementare le proprietà tribologiche di tale acciaio
senza comprometterne eccessivamente la resistenza a
corrosione. In particolare si cercheranno condizioni operative di trattamento che garantiscano:
- la formazione di composti ad elevata durezza in zona
superficiale;
- il mantenimento di un elevato contenuto di cromo non
legato (come carburo o nitruro) che possa garantire elevata resistenza a corrosione dell’acciaio.
In bibliografia sono presenti alcuni studi che riportano le
modifiche strutturali subite dall’acciaio 17-4PH quando
15
Memorie
sottoposto a trattamenti di nitrurazione a diverse temperature [4]. In particolare si ha che:
i) per basse temperature di trattamento l’azoto diffonde
nella fase α dell’acciaio, determinandone la saturazione e rendendo più stabile una struttura c.f.c. (vista la
sua azione γ-gena); in queste condizioni è possibile ottenere una fase α’N (martensite contenente azoto sovrasatura).
ii) già a temperature di 350°C è possibile la formazione di
una fase S metastabile, costituita da austenite espansa
(γN).
iii)per temperature pari o superiori ai 420°C la fase metastabile S scompare, a causa della formazione di nitruri
di cromo (trasformazione fase-S → α’N + CrN).
iv)per temperature superiori ai 450°C si ha completa
scomparsa anche della fase α’N a favore dei nitruri di
cromo (trasformazione α’N → α + CrN).
Dal punto di vista esclusivamente tribologico, la comparsa
di CrN determina un importante miglioramento delle proprietà superficiali dell’acciaio, innalzando la microdurezza
fino a valori superiori ai 1250 HV.
Nel presente lavoro si ricerca tuttavia una condizione di
trattamento che possa garantire sì un incremento della
durezza superficiale, ma senza compromettere eccessivamente la resistenza a corrosione del materiale. Trattamenti che portano alla formazione di elevati tenori di CrN
saranno pertanto da escludere, in quanto determineranno
un impoverimento di cromo tale da rendere l’acciaio non
più in grado di passivarsi.
Saranno pertanto possibili diverse strategie di trattamento:
- lavorare a temperature inferiori ai 450°C, cercando la
giusta combinazione di tempi e temperature di trattamento tali da incrementare la microdurezza superficiale
senza portare ad eccessivo impoverimento di cromo della matrice;
- lavorare nelle condizioni che garantiscono la maggior
microdurezza superficiale (450°C o superiori) tentando
di ripristinare la resistenza a corrosione mediante un
post-trattamento di ossidazione.
Un’ulteriore variabile è costituita dal trattamento di invecchiamento artificiale. Gli acciai PH sono infatti quasi
sempre sottoposti a tale trattamento, che induce la precipitazione di intermetallici estremamente fini, allo scopo
di incrementarne la durezza (Precipitation Hardening). Il
trattamento di invecchiamento, che può variare a seconda
della composizione dell’acciaio, consiste solitamente in un
riscaldamento a temperature nel range dei 500-600°C,
mantenimento a tali temperature per un periodo di tempo
di circa 4 ore e successivo raffreddamento in aria.
Per valutare l’influenza che il trattamento di invecchiamento può avere sui successivi trattamenti termochimici
applicati, sono stati previsti trattamenti di nitrurazione (a
parità di condizioni) sia su provini invecchiati (selezionando come invecchiamento standard l’H10251) che su provini
allo stato solubilizzato.
invecchiamento alla T di 550° (± 8°C) per un tempo di 4 ore e
successivo raffreddamento in aria
1
16
Sigla
campione
Tipo di trattamento
A
solubilizz. + invecchiamento H1025
B
solubilizz. + nitrurazione 520° 12h + postossidazione a 470°
C
solubilizz. + invecchiamento H1025 +
nitrurazione 520° 12h + post-ossidazione
a 470°
D
E
solubilizz.+ invecchiamento H1025 +
nitrurazione 470° 4h + post-ossidazione a
470°
H
solubilizz.+ nitrurazione 400° 16h
I
solubilizz. + invecchiamento H1025 +
nitrurazione 400° 16h
L
solubilizzazione (stato di fornitura)
M
N
solubilizz. + nitrurazione 440° 16h
Tab. 1: codifica campioni e specifiche trattamenti
termici e termochimici applicati.
Tab. 1: samples identification, in term of thermal and
thermochemical treatments applied
In Tabella 1 sono riportati i trattamenti individuati per condurre la ricerca. Oltre al campione allo stato di fornitura,
cioè privo di trattamenti (campione L) sono state previste diverse combinazioni di trattamenti che prevedono
appunto: i) invecchiamento H1025, ii) nitrurazione ionica,
iii) post-ossidazione. In particolare la nitrurazione è stata
condotta introducendo in camera per primo l’idrogeno, gas
riducente in grado di agire sugli ossidi passivanti dell’acciaio inossidabile e garantire pertanto un’adeguata preparazione delle superfici dei provini, consentendo la giusta
diffusione di azoto durante il processo di nitrurazione.
PARTE SPERIMENTALE
Tutti i campioni allo stato di fornitura sono stati lucidati
superficialmente prima di eseguire i trattamenti. A seguito
del trattamento di invecchiamento H1025 i provini si ricoprono di una patina superficiale bluastra (ossidi iridescenti). Come prassi si è deciso di rimuovere meccanicamente
(rilucidatura) tale patina prima di eseguire gli ulteriori trattamenti previsti (vedi Tabella 1).
Per il solo campione allo stato di fornitura (L) è stata eseguita un’analisi chimica mediante quantometro allo scopo
di verificare se la composizione dell’acciaio è in linea con
quanto previsto dalla denominazione 17-4PH.
Per l’intero set di campioni sono state eseguite prove di
Trattamenti termici
Composition
Carbon
0.07 max.
Manganese 1.00 max.
Phosphorus 0.040 max.
Sulfur
0.030 max.
Silicon
1.00 max.
Chromium 15.00-17.50
Nickel
3.00-5.00
3.00-5.00
Copper
Carbonium
plus Tantalum 0.15-0.45
Elemento
% (in peso)
Carbonio
0.030
Manganese
0.642
Fosforo
0.024
Zolfo
0.012
Silicio
0.525
Cromo
15.42
Nichel
4.32
Rame
3.56
Tantalio
0.005
Niobio (colombio)
0.259
Tab. 2: composizione chimica tipica di un acciaio 174PH (colonna a sinistra) e composizione chimica di un
provino L ottenuta a seguito di analisi al quantometro
(colonna di destra).
Tab. 2: typical chemical composition of 17-4PH stainless
steel (left column) and chemical composition of sample L
obtained by quantometer analysis (right column)
microdurezza Vickers superficiale (carico applicato 1kg
forza) per valutare gli incrementi di durezza ottenuti a seguito dei diversi trattamenti. I valori di HV1 sono stati ottenuti come media di 10 misure eseguite in diverse zone
della superficie del provino.
Sono inoltre stati acquisiti profili di microdurezza, eseguendo indentazioni Vickers con carichi di 100 g forza
sulle sezioni dei campioni preventivamente spianate e lucidate (sequenza di carte abrasive e panni con sospensioni
diamantate).
La resistenza a corrosione è stata valutata, per ciascun
campione di Tabella 1, mediante due prove di corrosione
accelerata secondo normativa ASTM-G5, operando sul lato
lucido e variando, nelle due prove, l’ambiente corrosivo:
nella prima prova è stato utilizzata una soluzione di cloruro di sodio (NaCl) 3.5%m/m (che simula l’azione aggressiva
degli ioni cloruro tipica di un’acqua marina), nella seconda
è stata utilizzata una soluzione di acido solforico (H2SO4)
0.5M (che simula un ambiente acido tipico da condensa
in atmosfera industriale e che rappresenta l’ambiente tipico di prova per gli acciai inossidabili secondo normativa
ASTM-G5)
Di seguito vengono riportate brevemente le specifiche della prova di corrosione accelerata:
- area superficiale di campione analizzata: 1cm2;
- ambiente: soluzione di NaCl 3.5%m/m oppure soluzione
di H2SO4 0.5M;
- polarizzazione eseguita mediante il seguente ciclo:
a) polarizzazione catodica dal potenziale di riposo del
2
Nel caso della prova in H2SO4 i potenziali applicati sono diversi, per assicurarsi di raggiungere la completa transpassivazione
dell’acciaio durante la prova: a) polarizzazione fino a (Er – 0.2)V;
b) polarizzazione fino a (Er + 1.8)V
Sigla
campione
A
B
C
D
E
H
I
L
M
N
Tipo di trattamento
Durezza
superficiale
HV1
solubilizz. + invecchiamento
375 ± 2
H1025
solubilizz. + nitrurazione 520°
12h + post-ossidazione a
751 ± 113a
470°
H1025 + nitrurazione 520°
628 ± 193b
12h + post-oss. a 470°
477 ± 14
4h + post-ossidazione a 470°
solubilizz.+ invecchiamento
H1025 + nitrurazione 470° 4h 527 ± 23
+ post-oss. a 470°
solubilizz.+ nitrurazione 400°
555 ± 19
16h
H1025 + nitrurazione 400°
480 ± 17
16h
solubilizzazione (stato di
335 ± 3
fornitura)
16h + post-ossidazione a
870 ± 30
440°
860 ± 24
16h
Tab. 3: microdurezze superficiali HV1. (a) elevata
deviazione standard in quanto la superficie del provino
presenta due zone a diversa durezza (una che fornisce
valori di poco superiori ai 600, l’altra con valori
superiori a 800) (b) elevata deviazione standard in
quanto la superficie del provino presenta due zone a
diversa durezza (una che fornisce valori compresi fra
450 e 500, l’altra con valori di poco inferiori agli 800).
Tab. 3: surface HV1 microhardness. (a) high standard
deviation due to the presence, on the sample surface, of
two regions with different hardness (a region with values
of almost 600, another region with values higher than 800)
(b) high standard deviation due to the presence, on the
sample surface, of two regions with different hardness (a
region with values ranging between 450 and 500, another
region with values of almost 800).
campione (Er) fino al potenziale (Er - 0.4)V2;
b) polarizzazione catodica dal valore raggiunto precendemente, (Er - 0.4)V, fino al valore (Er + 1.6)V2;
- velocita di scansione applicata: 0.0004 V/s;
- potenziali misurati rispetto ad elettrodo di riferimento
Ag/AgCl/KCl(saturo);
Al termine di ciascuna prova di corrosione sono state ac17
Memorie
quisite immagini della zona sottoposta a polarizzazione
mediante stereomicroscopio ottico, al fine di valutare la
morfologia di corrosione e la severità dell’attacco subito
a livello qualitativo. Tale indagine è utile per capire se vi è
presenza di corrosione localizzata e se la localizzazione segue particolari punti di innesco (ad esempio bordi grano).
RISULTATI E DISCUSSIONE
Composizione chimica
In Tabella 2 sono mostrate la composizione tipica di un acciaio 17-4PH (colonna a sinistra) e la composizione riscontrata sul provino L analizzato al quantometro (colonna a
destra). I risultati ottenuti al quantometro confermano che
la composizione chimica dell’acciaio soddisfa le specifiche
dettate dalla designazione 17-4PH.
Microdurezza superficiale HV1
In Tabella 3 sono riportati i valori di microdurezza superficiale (HV1) con relativo errore (deviazione standard associata alla serie di misure eseguite).
Osservando i risultati è possibile notare che:
- il trattamento H1025, applicato sul campione solubilizzato (L) come unico trattamento incrementa la durezza
superficiale di circa 40 unità HV;
- i trattamenti D, E, H, I portano a valori di HV che, considerando le deviazioni standard, possono essere ritenuti
confrontabili (spaziano in un range centrato intorno ai 500
HV); il valore migliore si ottiene sul provino H, nitrurato a
bassa temperatura per tempi lunghi e senza aver precedentemente subito il trattamento di invecchiamento.
- i trattamenti B e C, che prevedono l’applicazione della
maggior T di nitrurazione, portano ai valori di HV maggiori
ma sono caratterizzati da una spiccata disomogeneità del
trattamento; si ottengono infatti deviazioni standard molto
elevate motivate dal fatto che entrambi i provini presentano
zone ad elevata durezza (800 HV) e zone a durezza più contenuta. Probabilmente queste sono le condizioni ottimali
per massimizzare l’indurimento superficiale derivante dalla nitrurazione (formazione di CrN in elevata quantità) ma
sembra che l’acciaio inossidabile non sia in grado di subire
questo trattamento in maniera omogenea: il trattamento di
disossidazione della superficie applicato prima di introdurre in camera l’azoto (quindi in presenza di solo idrogeno)
non sembra pertanto adeguato per garantire omogeneità di
trattamento quando si nitrura ad elevata temperatura.
- i trattamenti M ed N, eseguiti a temperature intermedie
ma per tempi elevati, sembrano essere i più performanti,
portando a valori di microdurezza superficiale confrontabili con quelli ottenuti con i trattamenti a più elevata temperatura, ma assicurando una maggiore omogeneità di
trattamento.
18
- il trattamento di invecchiamento H1025, quando seguito
da trattamenti termochimici di nitrurazione, sembra ostacolarne l’efficacia. Confrontando le coppie di provini che
hanno subito gli stessi trattamenti ma che differiscono per
l’applicazione o meno dell’invecchiamento H1025 (B con
C, D con E ed H con I) si nota come le durezze superficiali
raggiunte siano sempre maggiori in assenza di invecchiamento (unica eccezione la coppia D con E dove i valori
sono comunque molto simili). Questo risultato fa supporre
che i fenomeni attivati durante l’invecchiamento influenzino il successivo trattamento di nitrurazione; evidenze a
riguardo emergono in letteratura [8] anche se non sono
riportati modelli esaustivi che spieghino il fenomeno.
Profili di microdurezza
In Figura 1 sono mostrati i profili di microdurezza HV0.1 ottenuti sulla sezione lucidata dei campioni trattati secondo
Tabella 1. I risultati confermano le evidenze ottenute dalle
microdurezze superficiali, in quanto:
- gli unici trattamenti in grado di determinare elevati valori
di durezza e una discreta profondità di indurimento sono
quelli eseguiti alle temperature maggiori (B) oppure a temperature di 440°C per tempi estremamente lunghi (M ed
N);
- il trattamento di invecchiamento H1025, quando seguito
da trattamenti termochimici di nitrurazione, sembra ostacolarne l’efficacia, mentre in assenza di post-trattamenti
di nitrurazione comporta un incremento della durezza del
materiale di oltre 50 punti HV (vedi confronto L-A).
- i trattamenti eseguiti a bassa temperatura (400°C, H ed
I) non modificano in maniera sostanziale la durezza degli
strati superficiali del materiale, risultando di fatto poco efficaci per gli scopi previsti.
Prove di corrosione
Provini A ed L
In Figura 2 sono mostrate le curve di polarizzazione ottenute sui provini A ed L in ambiente salino (NaCl 3.5% in
peso). Le curve mostrano un potenziale di pitting notevolmente differente: circa 0.15 V per il provino L e circa -0.05
V per il provino A. Il processo di invecchiamento sembra
quindi compromettere la resistenza a pitting in ambiente
clorurato, abbassando il potenziale di pitting dell’acciaio di
quasi 200mV. Il peggioramento è confermato anche dalla
corrente media nell’intervallo di passivazione che, per il
provino A (curva blu in figura) è leggermente superiore.
In Figura 3 sono mostrate le curve di polarizzazione ottenute sui provini A ed L in ambiente acido (H2SO4 0.5M). In
questo ambiente il trattamento di invecchiamento sembra
non influenzare minimamente la resistenza a corrosione
dell’acciaio, e le due curve appaiono (pur nella loro complessità, cioè con almeno due picchi di attivazione ben deLa Metallurgia Italiana - n. 4/2015
Trattamenti termici
Fig. 1 - Profili di microdurezza
HV0.1 eseguiti sulla sezione
lucidata dei campioni trattati
secondo Tabella 1.
Fig. 1 - HV0.1 microhardness profile
obtained from polished cross section
of samples of Table
Fig. 2 - Curve di polarizzazione
ottenute in NaCl 3.5%
Fig. 2 - Polarization curves obtained in
NaCl 3.5% solution
19
Memorie
lineati) perfettamente sovrapponibili.
I dati facilmente estrapolabili dalle curve di Figura 3, e che
permettono di delineare la resistenza a corrosione dell’acciaio 17-4PH in questo ambiente, sono:
- il potenziale di transpassivazione (circa 0.95 V per entrambi i provini);
- la corrente media nell’intervallo di passivazione (varia nel
range 5÷7x10-6 Acm-2 per entrambi i provini);
- la corrente massima raggiunta dal principale picco di attivazione (circa 6x10-4 Acm-2 per entrambi i provini);
I risultati emersi dall’analisi dei provini A ed L mostrano
pertanto che il trattamento di invecchiamento H1025 determina una diminuzione della resistenza a pitting indotto
da cloruri; il risultato può essere interpretato considerando una precipitazione di carburi di cromo a seguito del
trattamento termico (nonostante la presenza di elevato
tenore di elementi stabilizzanti quali niobio); questo peggioramento della resistenza a corrosione a seguito dell’invecchiamento (trattamento eseguito prevalentemente per
incrementare le proprietà meccaniche) valorizza il fatto di
intervenire mediante successivi trattamenti termochimici
superficiali che possano migliorare la resistenza a pitting
(nitrurazione ma soprattutto post-ossidazione).
Provini B, C, D, E
In Figura 4 sono mostrate le curve di polarizzazione ottenute sul lato lucidato dei provini B, C, D, E in ambiente salino (NaCl 3.5% in peso); nello stesso grafico sono riportate
le curve ottenute nel medesimo ambiente sui provini analizzati in precedenza: A ed L (non sottoposti a trattamenti
termochimici). I provini B, C, D, E hanno subito un trattamento di nitrurazione a temperature tipicamente utilizzate
per massimizzare i risultati in termini di resistenza ad usura; al trattamento è seguito una post-ossidazione.
Purtroppo è piuttosto evidente dai grafici che per tutti e
quattro questi provini l’intervallo di passivazione viene
quasi a scomparire. Osservando le curve nei tratti iniziali
(laddove i campioni A ed L mostravano un plateaux di passivazione a valori di corrente prossimi ai 10-6 Acm2) si vede
come i provini A, B, C, D evidenzino solamente un accenno
di plateaux di passivazione, a correnti molto più elevate
(circa 5x10-6 Acm-2) e di lunghezza (in termini di potenziale) estremamente ridotta. Dal punto di vista dell’intervallo
di passivazione i quattro campioni si classificano in questo
modo:
- campioni B e D (curve verde e fucsia): mostrano un intervallo di passivazione di lunghezza compresa fra i 100
e i 150mV, emergendo quindi come leggermente migliori
rispetto agli altri due;
- campioni C ed E (curve nera e azzurra): mostrano un intervallo di passivazione quasi inesistente (meno di 50mV
di ampiezza, paragonabile quasi ad un flesso);
Osservando attentamente le curve è possibile notare che,
a valori di potenziali maggiori (cioè una volta che il feno20
Fig. 3 - Curve di polarizzazione ottenute in H2SO4 0.5M
Fig. 3 - Polarization curves obtained in H2SO4 0.5M solution
Fig. 4 - Curve di polarizzazione ottenute in NaCl 3.5%
sui provini A, B, C, D, E, L
Fig. 4 - Polarization curves obtained from samples A, B, C,
D, E, L in NaCl 3.5% solution
meno di pitting è iniziato) tutte e quattro presentano una
specie di picco dopo il quale la corrente si abbassa notevolmente. Il fenomeno è molto evidente per la curva fucsia
(D), mentre è appena accennato per la curva azzurra (E).
Questo fenomeno è molto interessante è può essere ricondotto ad un arresto del processo di pitting una volta
che esso procede oltre la superficie dell’acciaio, ad opera
dell’azoto presente in soluzione solida nell’acciaio. L’azoto
presente nell’acciaio e liberato durante il processo di corrosione è in grado di reagire con gli ioni H+ e di abbattere
l’acidità all’interno dei pit che si stanno formando:
[N] + 4H+ + 3e- → NH4+ [9-11].
L’azoto agisce anche da stabilizzante per il film di passivazione dell’acciaio, rendendolo più resistente all’attacco degli ioni cloruro e può produrre ioni nitrato che aumentano
Trattamenti termici
la resistenza al pitting [12-14].
Il provino che sembra beneficiare maggiormente di tale
fenomeno protettivo è il D (curva fucsia) che sembra mostrare una vera e propria seconda zona di passivazione,
seppur molto limitata.
Considerando entrambi i fenomeni (plateaux di passivazione iniziale e ‘arresto del pit’ dovuto all’azoto in soluzione
solida) il campione D emerge come il più performante della serie dal punto di vista della resistenza a corrosione,
anche se rispetto ai provini non trattati (A ed L) questa proprietà sembra essere stata notevolmente compromessa.
Ad ogni modo dall’analisi emerge che le migliori performance in termini di resistenza a corrosione sono state
ottenute operando a T inferiori (470°C anziché 520°C) e
senza invecchiare il materiale.
Osservando il comportamento in acido solforico, Figura 5,
queste indicazioni vengono ulteriormente confermate:
- il provino C (invecchiato e trattato ad alta T, 520°C)
emerge come il peggiore della serie; oltre ad avere valori di corrente media di passivazione molto elevati questo
provino presenta un restringimento dell’intervallo di passivazione notevole;
- confrontando i provini non invecchiati (D con B, curve
fucsia e verde) si nota bene come la resistenza a corrosione peggiori con il trattamento ad alta T (520°C, provino B
curva verde); il peggioramento si ha soprattutto in termini
di corrente sul picco di attivazione e di corrente media di
passivazione;
- lo stesso risultato emerge confrontando i provini invecchiati (E con C, curve azzurra e nera): ancora una volta la
resistenza a corrosione peggiora con il trattamento ad alta
T (520°C, provino C curva nera);
Analizzando i risultati ottenuti fino a questo punto emerge
in modo chiaro come la T dei trattamenti influisca fortemente sulla resistenza a corrosione degli acciai trattati. Le
alte temperature compromettono la resistenza a corrosione sia in ambiente clorurato che in ambiente acido; questo
accade sia per il trattamento di invecchiamento (che essendo condotto a T di 550°C determina un peggioramento della resistenza a corrosione) che per il trattamento di
nitrurazione (i risultati ottenuti a 470° sono da preferire a
quelli ottenuti a 520°C).
Provini H ed I
In Figura 6 sono mostrate le curve di polarizzazione ottenute dai provini H ed I in ambiente salino (NaCl 3.5% in
peso); nello stesso grafico sono riportate le curve ottenute
nel medesimo ambiente su tutti i provini analizzati in precedenza.
Dal grafico emerge in maniera chiara che:
- entrambi i provini (H ed I) presentano una zona di passivazione iniziale (plateaux) più ampia rispetto a tutti i provini
trattati visti in precedenza (B, C, D E), ma ancora inferiore,
e a correnti maggiori, rispetto a quella dei provini non trattati (L ed A). Questo risultato è comunque molto importan-
Fig. 5 – Curve di polarizzazione ottenute in H2SO4
0.5M sui provini A, B, C, D, E, L
D, E, L in H2SO4 0.5M solution
sui provini A, B, C, D, E, H, I, L
D, E, H, I, L in NaCl 3.5% solution
te in quanto mostra come un trattamento di nitrurazione
eseguito a bassa T possa garantire una migliore preservazione della resistenza a corrosione mostrata inizialmente
dall’acciaio. Da questo punto di vista i due provini H ed I
sembrano comportarsi in modo molto simile fra di loro.
- entrambi i provini (H ed I) presentano una notevole
estensione della seconda zona di passività individuata nei
provini precedenti ed attribuita all’arresto del processo di
pitting a causa dell’effetto benefico dell’azoto presente in
soluzione solida nell’acciaio. Questo effetto è particolarmente evidente sul provino H e si estende fino a potenziali
prossimi a quelli di transpassivazione originali del provino
non trattato L.
In Figura 7 sono mostrate le curve di polarizzazione otte21
Memorie
nute dai provini H ed I in ambiente acido (H2SO4 0.5M);
nello stesso grafico sono riportate le curve ottenute nel
medesimo ambiente su tutti i provini analizzati in precedenza. Dal grafico è possibile notare che le curve relative
ai provini trattati a bassa T possiedono:
- un intervallo di passivazione esteso quanto quello del
provino non trattato (L);
- una corrente media nell’intervallo di passivazione più elevata di quella del provino non trattato, ma posizionata a
valori minimi fra quelli riscontrati sui provini trattati. Per
vedere meglio questo comportamento è stato realizzato il
grafico di Figura 8 dove le curve caratterizzate dalle oscillazioni più ampie sono state trattate con un filtro passabasso per permettere di individuare meglio il valore di
corrente media di passivazione. Dal grafico si nota come
il provino I si collochi a valori minimi di corrente media
(paragonabili a quelli del provino D visto in precedenza)
mentre il provino H a valori lievemente superiori.
Fig. 7 - Curve di polarizzazione ottenute in H2SO4
0.5M sui provini A, B, C, D, E, H, I, L
D, E, H, I, L in H2SO4 0.5M solution
- assenza quasi completa del primo picco di attivazione;
da questo punto di vista i provini H ed I sembrano comportarsi meglio rispetto a tutti i provini analizzati, compresi i
non trattati.
Provini M ed N
In Figura 9 sono mostrate le curve di polarizzazione ottenute dai provini M ed N in ambiente salino (NaCl 3.5% in
peso); nello stesso grafico sono riportate le curve ottenute
nel medesimo ambiente su i campioni non nitrurati (A ed
L) e sui provini che in precedenza avevano mostrato i migliori risultati in tale ambiente (H ed I).
Dal grafico emerge che:
- entrambi i provini (M ed N) presentano una zona di passivazione iniziale più ampia e a correnti inferiori rispetto ai
migliori risultati ottenuti in precedenza (H ed I) e, per la prima volta, si hanno curve confrontabili con il campione non
nitrurato (ma invecchiato) A. Questi risultati sono i migliori
ottenuti sull’intera serie di trattamenti testati, e possono
essere messi in relazione con l’ottenimento di una microstruttura che, in accordo con i dati bibliografici [4] presenta un elevato tenore di azoto disciolto (struttura α’N).
- minima invece la differenza fra i due provini M ed N, risultato che consente di affermare che, ottenendo una nitrurazione performante in termini di resistenza a corrosione,
non è necessario applicare un post-trattamento di ossidazione.
In Figura 10 sono mostrate le curve di polarizzazione ottenute dai provini M ed N in ambiente acido (H2SO4 0.5M);
nello stesso grafico sono riportate le curve ottenute nel
medesimo ambiente su tutti i provini analizzati in precedenza. Dal grafico è possibile notare che le curve relative
ai provini M ed N non risultino particolarmente performanti in ambiente acido. Mentre per il provino M tale com22
Fig. 8 - Curve di polarizzazione ottenute in H2SO4 0.5M
sui provini A, B, C, D, E, H, I, L (smoothing applicato
alle curve con oscillazioni più ampie utilizzando un
filtro passa-basso)
Fig. 8 - Polarization curves obtained from samples A, B,
C, D, E, H, I, L in H2SO4 0.5M solution (the curves that
showed large current oscillations were smoothed with a
low-pass filter)
portamento è probabilmente attribuibile ad un incremento
dell’area superficiale attiva a causa del post-trattamento
di ossidazione, per il provino N si hanno risultati simili a
quelli ottenuti per il provino D (trattato a 470°C).
Questi risultati permettono di concludere che le ottime
performance di resistenza a corrosione in ambiente clorurato ottenute con i trattamenti M ed N derivano da una
struttura ad elevato contenuto di azoto disciolto, e non
dall’assenza di formazione di precipitati CrN.
L’azoto disciolto, come visto in precedenza, è in grado
di rendere meno severi gli attacchi di tipo caverniforme
(pitting) [9-14], mentre in un ambiente decisamente acido
Trattamenti termici
CONCLUSIONI
L’attività di ricerca ha permesso di caratterizzare dal punto di vista della resistenza a corrosione in due differenti
ambienti (clorurato ed acido) diverse combinazioni di trattamenti termici e termochimici applicati all’acciaio 17-4PH
allo scopo di migliorarne la resistenza meccanica e le proprietà tribologiche. I risultati più importanti emersi dalla
ricerca sono di seguito riportati:
sui provini A, H, I, L, M, N
Fig. 9 - Polarization curves obtained from samples A, H, I,
L, M, N in NaCl 3.5% solution
1) i trattamenti che permettono di ottenere il miglior compromesso fra durezza superficiale, profondità di indurimento e resistenza a corrosione sono quelli condotti per
tempi lunghi (16 ore) alla temperatura di 440°C; in tali
condizioni è possibile ottenere un elevato tenore di azoto disciolto e una minima formazione di nitruri di cromo,
risultato che porta ad un’ottima resistenza a corrosione
in ambiente clorurato e ad una notevole microdurezza superficiale (superiore ad 800HV). In tali condizioni non si
evita comunque l’impoverimento di cromo della matrice
dato dalla precipitazione dei CrN, fenomeno che porta ad
avere una resistenza a corrosione in ambiente acido discretamente compromessa rispetto a trattamenti analoghi
eseguiti a più bassa temperatura (400°C).
2) i trattamenti a bassa temperatura (400°C) permettono
di mantenere ottime proprietà di resistenza a corrosione
sia in ambiente acido che clorurato, ma determinano un
incremento della durezza superficiale molto modesto.
Fig. 10 - Curve di polarizzazione ottenute in H2SO4
0.5M sui provini A, B, C, D, E, H, I, L, M, N
Fig. 10 - Polarization curves obtained from samples A, H, I,
L, M, N in H2SO4 0.5M solution
(H2SO4) risulta fondamentale il contenuto di cromo disponibile in lega (non legato sotto forma di carburi e nitruri).
I campioni H ed I, ottenuti a bassa temperatura, rispondono bene a corrosione anche in ambiente acido in quanto
risulteranno privi di precipitati di CrN (come confermato
anche dalle scarse proprietà di indurimento superficiale
riscontrate dalle misure di HV superficiale ed in sezione),
mentre i campioni M ed N rispondono bene a corrosione
solo in ambiente clorurato in quanto, essendo stati trattati
per tempi lunghi a T superiori (440°C) mantengono sì un
elevato contenuto di N disciolto, ma avranno sicuramente
dato luogo anche alla precipitazione di nitruri di cromo (fenomeno ancora una volta confermato dagli ottimi valori di
HV riscontrati per questi provini).
3) il trattamento di invecchiamento H1025, se applicato
prima di eseguire la nitrurazione, influenza negativamente i risultati ottenibili dal trattamento termochimico. L’incremento di durezza a cuore ottenibile con il trattamento
H1025 è comunque perseguibile anche durante il normale
processo di nitrurazione (che opera a temperature tali da
consentire i fenomeni di precipitation hardening), pertanto
risulta conveniente trattare termochimicamente provini di
acciaio 17-4PH allo stato solubilizzato.
BIBLIOGRAFIA
[1]
[2]
[3]
[4]
[5]
[6]
[7]
[8]
[9]
[10]
[11]
[12]
[13]
[14]
W.T. Chien, C.S. Tsai, J. Mater. Proc. Technol. 140 (2003) 340.
P. Li, Q. Cai, B. Wei, X. Zhang, J. Iron. St. Res. 13 (2006) 73.
P. Kochmanski, J. Nowacki, Surf. Coat. Technol. 200 (2006)
6558.
G. Li, J. Wang, C. Li, Q. Peng, J. Gao, B. Shen, Nuclear Instruments
and Methods in Physics Research B 266 (2008) 1964.
J. Wang, H. Zou, C. Li, Y. Peng, S. Qiu, B. Shen, Nucl. Eng. Design
236 (2006) 2531.
J. Wang, H. Zou, C. Li, S. Qiu, B. Shen, Mater. Charact. 57 (2006)
274.
F. Christien, R. Le Gall, G. Saindrenan, Scripta Mater. 48 (2003)
11.
P. Kochmański, J. Nowacki, Surface & Coatings Technology 202
(2008) 4834.
C.X. Li , T. Bell, Corrosion Science 46 (2004) 1527.
S.D. Chyou, H.C. Shih, Corrosion 47 (1991) 31.
H.J. Grabke, The role of nitrogen in the corrosion of iron and steels, ISIJ International 36 (1996) 777.
I. Olefjord, L. Wegrelius, Corrosion Science 38 (1996) 1203.
H. Baba, T. Kodama, Y. Katada, Corrosion Science 44 (2002)
2393.
U. Kamachi Mudali, P. Shankar, S. Ningshen, R.K. Dayal, H.S. Khatak, B. Raj, Corrosion Science 44 (2002) 2183.
23
Memorie
Nitriding and post-oxidation treatments
on 17-4ph stainless steel: optimization of the
process to preserve high corrosion resistance
Keywords: Stainless Steel - Corrosion - Thermal Treatment - Material Characterization
17-4PH stainless steel is usually heat treated to increase its mechanical properties. In order to obtain a further improvement of the wear resistance of this steel, it is possible to apply thermochemical treatments, such as nitriding.
The untreated 17-4 PH stainless steel has a remarkable corrosion resistance conferred by the high chromium of this
alloy, so it’s important to evaluate how the application of a thermal or thermochemical treatment can affect this
property.
The aim of this work is to check how the application of thermochemical treatments, usually used to increase the mechanical and wear properties of iron alloys, affect the corrosion resistance of this steel and to optimize the process
variables (considering also the possibility to add subsequent treatments, such as post-oxidation) in order to identify
the best treatment conditions in order to preserve a good corrosion resistance.
For this purpose accelerated corrosion tests were performed, through the acquisition of polarization curves in an
electrochemical cell, on specimens subjected to different combinations of heat and thermochemical treatments
(H1025 aging, nitriding, post-oxidation). In addition to the electrochemical characterization, the specimens were
characterized by surface HV microhardness tests and by HV microhardness profiles along their cross-section, to
assess the effectiveness of the applied treatments in terms of mechanical properties and wear resistance.
Despite the best results (in terms of increasing of surface hardness and depth of hardening) have been achieved
with treatments that significantly compromise the corrosion resistance of the steel, the work has allowed to identify
and optimize sequences of treatments that preserve almost completely the corrosion resistance of the steel, while
increasing the surface hardness up to values of 850HV.
24
Nanomateriali
Validity of Wulff construction used for
size-dependent melting point of nanoparticles
S. Zhang, L. Zhang, L. Chen
An integrated model based on the variant of Ba/Bt, is established to predict size-dependent melting point of
nanoparticles by considering the geometric and energetic characteristics of Wulff construction. Ba is the rest
bond number and Bt denotes the total bond number without broken bonds in a Wulff construction. Without
any adjustable parameters, this model predicts a decreasing trend of melting point with the size dropping for
nanoparticles. The good agreement between theoretical predictions and the evidences in experiments and
molecular dynamic simulation confirms the validity of Wulff construction in describing thermodynamic behaviors
of nanoparticles even with no need in considering their crystalline structures.
Keywords: Metals - Nanoparticles - Wulff construction - Melting
Introduction
The thermodynamic behavior of nanocrystals differs from
that of the corresponding bulk materials mainly due to the
large value of surface-to-volume ratio, which strongly influences both the chemical and physical properties in comparison with the bulk counterpart [1-4]. This is because the
surface/volume ratio depends on both size and shape, and
the size and shape or structure strongly influences many
fundamental properties of nanoparticles [5]. However, the
shape or structure is strongly depending on size of materials
[6-8]. It has been predicted that Na [9] and Mo [10] substances with a bulk bcc structure would have fcc or more like icosahedron structures for nanoparticles. This is because the
fcc or icosahedron structures are more compact than the
bcc structure and provide a lower surface energy than the
S. Zhang, L. Zhang
School of Mechanical Science and Engineering,
Jilin University,
Changchun 130025, China
[email protected]
S. Zhang
Department of Materials Science and Engineering,
Jilin JianZhu University,
L. Chen
Department of Municipal and environmental
engineering, Jilin JianZhu University,
[email protected]
bcc one. It is also found that Co nanoparticles with radius
below 10 nm prefer to form a fcc structure, rather than bulk
hcp one [11]. Moreover, many other nanoparticles bound by
van der Waals or metallic forces (such as Mg, Ca, Sr, Ni and
Ba) exhibit structures with fivefold axes of symmetry, i.e.,
icosahedron structure, despite the fact that the bulk metals
exhibit hcp, fcc or bcc packing [12]. It should be noted that
nanoparticles must display the bulk crystalline structure at
larger r (r shows the size of nanoparticles). Therefore, we
can expect that it is the surface energy controlling the shape or structure of nanoparticles, namely, the structure of
nanoparticles is the one with the smallest surface energy.
Since the shape or structure affect most properties of nanoparticles, it is necessary to be investigated. It is known that
Wulff construction, which is developed by minimizing surface energy for a given enclosed volume, is the standard method for determining the equilibrium shape of crystals at the
macroscale level. This requirement for small surface energy
is also applied to nanoparticles, since the surface energy is
the major energy contribution for them. So, considering the
compact packing of nanoparticles mentioned above, the geometrical and energetic properties of Wulff construction for
fcc crystal is taken for describing nanoparticles in this work.
Through introducing a variant Ba/Bt to describe the geometric characteristics of Wulff construction, a model without
any adjustable parameter is obtained to estimate size-dependent melting point of nanoparticles, where Ba is the rest
bond number and Bt denotes the total bond number without
broken bonds in a system. The good agreement between
our model predictions and experimental results suggests
that it is valid that taking Wulff construction as nanoparticles’ structure for predicting their melting temperature.
25
Memorie
Model
It is known, cohesive energy that describes the bond
strength directly, is an effective variable to determine the
thermal stability of nanocrystals. With the size reduction,
the decline of melting point is an obvious, which implies
the lowered thermal stability of nanocrystals. In fact, there
is an empirical correlativity between E0 and Tm0 functions,
by defining E0 and Tm0 as bulk cohesive energy and bulk
melting point [13, 14],
(1.1)
In Eq. (1.1), kb is the Boltzmann’s constant. According to
Eq. (1.1), if applying this relation to the nanoscale a similar
treatment for the relationship between E(r) and Tm(r) functions can be expected as a first approximation, that is,
(1.2)
Therefore, combining Eqs (1.1) and (1.2), the ratio of the
melting temperature of the nanoparticles versus that of
the bulk can be read as,
(2)
For a system, E(r) function has been derived by introducing
the variant of Ba/Bt, that is [15, 16],
same atom number. Arriving here, a Wulff construction is
established, and the size or diameter of a Wulff construction can be altered by controlling the the atom number
on edge. To obtain the Ba/Bt of a Wulff contruction, the
he total atoms number (Nt) and surface atoms number
(Ns) must be known. Let n denoting the atom number on
a edge, Nt and Ns can be resolved, e atom number on a
edge, Nt and Ns can be resolved,
(4.1)
The number of surface atoms can be expressed as following:
(4.2)
In fact, the value of Ns includes the number of the atoms
on (111) facets (N111), the number of atoms on (100) faces
(N100), the number of atoms at edges (Ne) and the number
of atoms on vertex Nv, that is Ns = N111+N100+Ne+Nv. From
mathematic point, N111 = 8(3n2-9n+7), N100 = 6(n-2)2, Ne =
36(n-2) and Nv = 24. In addition, the coordination number
should be resolved to obtain Ba/Bt. However, the coordination number for atoms at different sites is also different,
that is Z111 = 9, Z100 = 8, Ze = 7 and Zv = 6, respectively. So,
the average coordination number of surface atoms can be
expressed:
(5)
Then, one can obtain Ba = ZsNs/2. It is also easy to get Bt
value, since Bt = NtZb/2. Nt is given by Eq. (1). Zb is the
coordination number of bulk interior atoms, and Zb = 12
for fcc structure. Therefore, Ba/Bt can be expressed as the
following formulation,
(3)
The broken bonds of the atoms on surfaces inevitably lead
to the instability of materials in nano-scale. Thus, as long
as Ba/Bt is known, E(r) or Tm(r) is obtained. However, it is
necessary to know nanopaticle’s structure and size, since
both of them decide Ba and Bt values. For most metallic
nanoparticles, the most proper structure of nanoparticles
could be Wulff construction.
It is clear that Ba or Bt is strongly dependent on the size
and shape, since Ba and Bt actually are the multiplying results between the atom number and the average coordination number [15], namely Ba/Bt = ZsNs/ZbNt, where Zs and
Zb are average coordination number for surface atoms and
bulk interior, and Ns, Nt are the number of surface atoms
and total atoms in a system, respectively. So Eq. (3) indicates the size and shape dependences of cohesive energy,
and even for melting point of nanoparticles.
Wulff construction is a segment of fcc (faced-centeredcubic) crystal. By truncating a octahedron, one can obtain
a polyhedron with fourteen facets. There have six square
(100) facets and eight hexagonal (111) facets at its surface,
in which three edges of the hexagon are in common with
square (100) facets, while the remaining three edges in
common with hexagonal (111) facets. And each edge has
26
(6)
It is clear that n is related with the size of Wulff construction. So the value of Ba/Bt for Wulff construction is relying
on size. It is clear that different shape has different Ba/Bt
value. That is to say the value of Ba/Bt is simultaneously
related with both size and shape. Assuming the radius D
of Wulff construction as the biggest distance from center
atoms to surface atoms, D has the following expression,
(7)
with h being atomic distance.
Substituting Eqs. (3) and (6) into Eq. (2), size-dependent
melting point can be expressed
(8)
Results and discussions
Fig. 1 shows the comparison between model predictions in
light of Eq. (8) and experimental results for melting points
Nanomateriali
of several metallic nanoparticles. It is clear that a good
agreement between them is found. As expected, Tm(r) is a
continuous function of r and decreases monotonically as
r decreases, leading to the lowered thermal stability. This
is because of the decreased Ba/Bt value. The results displayed in Fig. 1 confirm the success of Wulff construction
for describing the geometric and energetic characteristics
of nanoparticles almost throughout the whole size range.
This is because the variant Ba/Bt appearing in Eq. (3) is related not only to size, but also to shape or structure. It can
effectively change the Ec(r) and Tm(r) value by swaying the
thermodynamic stability due to the change in Bt value. As
r decreasing, Ns relatively increases, which results in the
decrease in the total bond number and the increase in the
broken bond number. Moreover, it should be noted that Eq.
(8) is still valid for In and Sn nanoparticles with their bulk
structures being tetragonal. Therefore, it is expected that
taking Wulff construction as a standard shape to describe
nanoparticles is reasonable in full size range from micro to
macro without taking structure change into account.
For larger particles with r > 10 nm, the validity of Wulff
construction is clear. This is because, the change in bond
energy in comparison with that in bulk interior, is small,
and Ba = Bt for larger particles. However, the assumption
used in Eq. (10) also results small difference for smaller
particles, since only surface bond relaxation is considered in Eq. (3). In fact, except surface atoms, interior atoms
also become unstable compared to bulk interior, resulting
in larger estimation of Eq. (8). In addition, the defect or
vacancy in a nanoparticle is not considered in this work,
which means the result of ideal crystal by using Eq. (8).
This also may lead to small overestimation of Eq. (8), of
necessity for small particles, as presented in Fig. 1. Despite the existing errors, Eq. (8) can still be regarded as
a valid and simple way to predict Tm(r) values even in full
size range. It should be note that for small nanoparticles
with r < 5 nm, the validity of Eq. (8) implies that the nanoparticles possess close-packed structure whatever the
bulk structure is. Based on Eq. (8), to determine the Tm(r)
or E(r) values of nanoparticles, there is no need to know
surface energies or shape, and even other thermodynamic
information but atomic distance and the size of nanoparticles.
In the previous studies of nanoparticles, a spherical shape
is usually taken into account, and the reasonability of this
action for melting point of nanoparticles is also presented
by taking the ratio of surface/volume as the only variant
[17-20]. To further confirm the validity of Wulff construction developed in this work, the ratio of surface to volume
(δ) may explain this point. For spherical nanoparticle, can
be simply determined as,
(9)
For comparison, δ function for Wulff construction is given,
(10)
Fig. 1 - The comparison of Tm(r) functions between
model prediction in terms of Eq. (7) (solid lines) with
the help of Eq. (6) and experimental results of Al, Pb,
In and Sn nanoparticles, respectively, where ●, ◄, ♦,
and ► show experimental results [6]. The h used in
Eq. (6) are separately 0.3164 nm, 0.3870 nm, 0.3684
nm, and 0.3724 nm for Al, Pb, In and Sn elements.
Fig. 2 - The comparison of δ(r) functions between
sphere (dotted line) and Wulff construction (solid line)
in light of Eqs. (8) and (9) respectively, where h = 0.3
nm is taken for simplicity.
To further describe the validity of Wulff construction in describing the shape of nanoparticles, the comparison of δ by
Eqs (9) and (10) is made. As shown in Fig. 2, the changes
in δ with respect to size are presented. Similar trend for
spherical nanoparticles with that of Wulff construction is
found. As r increases, δ decreases and δ → 0 with r → ∞.
And the difference in δ between spherical shape and Wulff
construction decreases with r increasing. When r > 6 nm,
the difference between them is almost indistinguishable.
Note, with r decreasing, the particles is no longer a sphere
27
Memorie
one, thus the model based on sphere consideration is not
reasonable. However, Wulff construction is the truncated
octahedron with fcc structure and due to the small surface
energy, some small particles perfer to take this shape. As
a result, the model established in this work could be used
to predict the thermodynamic stability of small particles.
In addition, our results also strongly support the assumption of spherical shape usually considered for particles.
Conclusion
By utilizing the geometric characteristic of Wulff construction, the size-dependent melting point of nanoparticles is
modeled with the help of the variant Ba/Bt. Similar to other
melting models, this model predicts the decreasing trend
of melting point when the size is dropping. This mainly
arises from the lowered bond number in a nanoparticle
if compared with its bulk material. The consistency of the
model predictions and experimental results suggests the
validity of Wulff construction, one hand to describe the
shape or structure, and the other to describe the thermodynamic stability of nanoparticles.
ACKNOWLEDGEMENTS
The authors acknowledge the financial supports of the
Science and Technology project of Jilin province education department during the Twelfth Five-year Plan Period
( No.2013232). National Natural Science Foundation of
China (grant No. 51101067), Natural Science Foundation of Anhui Higher Education Institutions of China (No.
KJ2012B159), Open Foundation of Key Laboratory of Automobile Materials of the Ministry of Educations, Jilin University (No. 12-450060481289).
28
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Acciai
Microstructural characterization and
production of high yield strength rebar
E. Mansutti, G. Luvarà, C. Fabbro, N. Redolfi
Various technical standards from all over the world set out the mechanical and chemical characteristics for high
yield strength rebar. High yield strength rebar - as defined in this study – is applied to all concrete reinforcement
steel grades which require a minimum yield strength of 600MPa. The standards concerning rebar production were
reviewed in order to select all the possible grades that come under the above-mentioned definition.
This research project aims to determine if by applying an in-line quenching and self-tempering process, the
technological requirements for high yield strength rebar, as specified in the standards, can be met, in order to
optimize the chemical composition and save on alloying elements. The work can be divided into two different
phases. The preliminary phase took place in the metallurgical laboratory of Danieli’s research center and the
second phase in an industrial plant. Tests done in the laboratory set out to evaluate the effect of quenching
and chemical composition on the rebar’s final mechanical properties and microstructure. The purpose of the
industrial-scale tests was to evaluate the potential of DANIELI’s in-line quenching and self-tempering process,
referred to as QTB (Quenching and Tempering Bar process), applied to high-strength steels. At the end of the
lab tests, three different chemical compositions were selected, deemed suitable for the production of high yield
strength rebar. In the industrial-scale tests it was then possible to evaluate the performance of the QTB process
in the production of high yield strength rebar in terms of operating flow rates / pressures, optimized chemical
compositions, productivity and process stability.
Keywords: Rebar - Yield Strenght – Quenching - Microstructure
INTRODUCTION
The application of high yield strength rebar is provided for
in various technical standards from all over the world, such
as the US, Russia, Korea and Japan.
Russia, for example, introduced the concept of high-yield
rebar (980MPa) back in 1982, which was then developed
further in GOST 10884 issued in 1994.
The mentioned GOST standard takes advantage of the
known effect of silicon on enhancing elastic limit, allowing
it to be added up to a maximum of 2.3% so that the steel
can be included in the At1200 class (or class VI, considering the former standard), which corresponds to a minimum yield strength of 1200MPa.
Less indicative, however, is the recent Korean standard
that for the SD700 class only specifies a limitation regarding equivalent carbon (CeqIIW = 0.63).
In Japan in 1993 a research project was carried out, referred to as “New RC Project”, which was then incorporated
into the National Building Code [1] [2].
The US has published the most recent ASTM standards
on this subject. Both standard A615/A615M and A706/
E. Mansutti, G. Luvarà, C. Fabbro, N. Redolfi
Danieli & C. Off Meccaniche, Buttrio
A706M introduced “Grade 80”, which not only requires minimum yield stress values but also particularly high minimum UTS values (725MPa for standard A615 and 690MPa
for A706). In addition, the A706 is more demanding in
terms of Rm/Rp ratio, maximum carbon content and Ceq;
in practice this makes it more complicated to apply on-line
heat treatments in rolling mills, requiring greater attention
to be placed on chemical composition. It is also important
to bear in mind that compared to European standards, US
standards are more stringent in terms of statistical reliability of technological values, requiring rebar producers to
guarantee yield strengths that are significantly higher than
the minimum requirements of the standard.
Again, in the US market standard A1035 provides for the
possibility of producing high-tensile corrosion-resistant rebar through high chrome content (around 9%) and by controlling the final microstructure by taking advantage of the
new technologies for in-line heat treatments [3].
It is interesting to note that in the US market various rebar
producers are pushing for the introduction of high yieldstrength grades (such as proposing classes “100” and
“125”), even if current market demand for this type of product is low.
In China there are no reference standards for equivalent
grades, although some studies refer to the use of V-N microalloyed steels and ultrafine grained steels [4][5][6].
29
Memorie
COUNTRY
Ref. Standard
Maximum yield strength
Remarks
RUSSIA
GOST 10884-94
1200 MPa
High yield strength with addition of silicon up to 2.3%
UKRAINE
DSTU 3760-06
1000 MPa
-
JAPAN
“New RC Project 1993”
980 MPa
Also includes grades @ 1275
MPa but only for transverse
reinforcement applications
USA
ASTM A1035-14
830 MPa
(120 ksi)
High yield strength by controlling microstructure
KOREA
KS D3504-11
700 MPa
Ceq increase allowed up to
0.63
ENGLAND
BS 6744-01+
A2:09
650 MPa
Stainless steel rebar
INDIA
IS 1786-08
600 MPa
Microalloyed steel with
maximum Ceq of 0.53
CHINA
GB1499.2-07
500 MPa
Ceq max 0.55
Tab. 1 – Overview of international standards for high-tensile rebar. Ceq as per IIW standard:
(C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5.
Fig. 1 – Diagram of a timed quenching system and example of a sample with a thermocouple attached to it (top).
Photo showing the quenching station with heating oven and timed quenching system.
30
Acciai
Fig. 2 – Temperature trend measured
at the core of a DIA 16mm rebar during
the test
Table 1 summarizes the main international reference standards for high yield strength rebar.
Civil engineering applications until now have been rather limited even if this type of rebar is promising as it simplifies
the reinforcement of concrete [7].
GOAL
This study aims to examine the possibility of meeting the
technological requirements specified in various international standards for high yield strength rebar, using the in-line
quenching and self-tempering process, thereby optimizing
the chemical composition with considerable savings in alloying elements.
Tests done in the laboratory set out to evaluate the effect
of quenching and chemical composition on the rebar’s final mechanical properties and microstructure.
The industrial-scale tests then evaluated the potential of
DANIELI’s in-line quenching and self-tempering process,
referred to as QTB (Quenching and Tempering Bar process), applied to high-strength steels.
- Air cooling it down to 850°C and then quenching it (from
1 to 5 seconds)
- Interrupting the cooling process by self-tempering of the
material surface, and final air cooling.
DIA 16mm rebars with three different compositions were
selected for the experiment (see Table 2) in compliance
with specific international standards.
The following three compositions were used:
- Composition#1 with medium carbon content and high
silicon content;
- Composition #2 with low carbon content and medium
Mn content;
- Composition #3 with high Mn content and medium Si
content.
The aim of the experimental plan was to determine the
combined effect of C, Mn and Si on hardenability and performance in terms of mechanical and microstructural properties. In particular the effect of a composition with lower
Mn and higher C and Si contents (such as composition
#3 for example), was compared to the other two chemical
compositions with lower carbon and higher Mn contents.
LABORATORY TESTS
LABORATORY EXPERIMENTAL PROCEDURES AND MATERIAL CHARACTERIZATION
In order to study the mechanical and microstructural properties of various rebars subjected to a quenching and
self-tempering process, a device (suitable for different
diameters) was set up to heat-treat samples of rebar that
were previously fitted with thermocouples (Figure 1).
The system is made up of:
- Heating oven with inert atmosphere (Ar)
- Brine quenching tank with timed immersion system
For the experiment it was decided to use a DIA 16 rebar
with a length of about 250mm.
Shown in Figure 2 is a representative trial heat cycle of a
sample subjected to testing.
The experiment involved:
- Heating the sample to 900 °C
- Soaking it for 5 minutes
Samples of each rebar composition shown in Table 2 were
quenched at increasing immersion times from 1 to 5 seconds while the core temperature was continuously monitored. For each test both microstructural properties and
mechanical strength were analyzed and measured.
To facilitate the comparison of results, before performing
the tests, various heat treatments were considered to decide which one would be used to determine the same prior
austenitic grain size for all the samples. This made it possible to use the same heat treatment (briefly described in
the previous chapter) for all three steel grades studied.
The quenched rebar pieces were subjected to a tensile
test to determine their mechanical properties.
Figure 3 shows a growing linear trend of yield strength up
to a maximum of 1000 MPa for compositions #2 e #3,
31
Memorie
Composition
%C
%Mn
%Si
%P
%S
%V
%Al
%Cu
%Cr
%Ni
N ppm
Ceq
#1
0.36
0.67
0.96
0.033
0.026
0.008
0.003
0.06
0.05
0.03
84
0.49
#2
0.22
0.98
0.18
0.018
0.015
0.003
0.004
0.28
0.10
0.08
117
0.41
#3
0.19
1.31
0.50
0.033
0.033
0.006
0.003
0.03
0.02
0.01
58
0.41
Tab. 2 – Result of the chemical analyses performed on samples from lots selected for the experiment;
elements not shown on the table are present only in trace amounts. Ceq according to standard IIW:
(C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5.
while composition #1 exceeds 1200 MPa.
On Figure 3 it can be noted that elongation diminishes
along a linear path as quenching time increases, with all
three steels following the same trend, while the decrease
in Rm/Rp ratio is less marked.
Because of its higher carbon and silicon contents, chemical composition #1 is able to reach the required yield
strength within a shorter quenching time. Moreover, even
with higher carbon and silicon contents, ultimate elongation is not penalized for up to 3 seconds of quenching
(which makes it possible to obtain a product with a yield
strength of 980MPa).
Composition #2, which has lower C, Mn and Si contents,
produces the lowest elongation value, even with relatively
short quenching times (1sec).
On next page is reported a summary of the results for:
- Analysis of microhardness (HV0.3) within the cross-sectional area of a rebar subjected to different cooling times;
- Description of the microstructure observed at the core of
the rebar with temperature measured at the end of immersion (thermocouple placed at the core of the rebar).
In general, a gradual increase is noted in the presence of
rapidly cooled structures at the core of the rebar, and even
completely hardened structures resulting from quenching
times of between 2.5 and 3.0 seconds.
Compared to composition #2, the increase in Mn and Si
for chemical composition #3 leads to a slight rise in martensite hardness at the end of the cycle (in both cases
cooling time was 4”).
Increased material hardenability due to higher Mn and Si,
together with the effect of tempering stability provided
by the silicon, still leads to increased hardness within the
cross-sectional area of a rebar quenched for the same
amount of time.
A comparison of the above results with those of chemical composition #1 show that the high C and Si contents,
which ensure greater hardenability, make it possible to
achieve complete hardening with shorter quenching times.
The result obtained with a shorter quenching time (2 sec)
and higher final temperatures (550 °C) is comparable to
the performances of the other steel grades.
The strategy of using higher amounts of carbon and silicon
while reducing the amount of manganese is only effective
if managed properly through controlled cooling.
Steel
2#
3#
1#
Cooling
time
[s]
1.0
2.0
2.5
3.0
3.5
3.5
4.0
4.0
1.0
2.0
2.5
3.0
3.5
4.0
2.0
3.0
3.0
4.0
4.0
5.0
Rp0.2
[MpA]
Rm
[MpA]
Rm/Rp
A%
A5d
%
420
600
861
880
920
920
1068
1000
434
660
650
920
1040
1114
650
980
980
1000
1120
1200
603
888
1053
1005
1118
1132
1296
1329
589
803
829
1025
1234
1282
830
1188
1276
1293
1527
1578
1.44
1.48
1.22
1.14
1.22
1.23
1.21
1.33
1.36
1.22
1.27
1.11
1.19
1.15
1.28
1.21
1.30
1.29
1,36
1.,32
16
6
10
8
5
6
3
4
17
12
12
8
9
7
15
10
8
4
7
5
16
3
9
7
4
5
3
3
20
3
14
10
9
4
16
11
9
3
5
5
Red n.: breakage outside calibrated lenght
Fig. 3 – Change in mechanical properties as quenching
time increases (right bottom). The measured
mechanical properties are reported on the top right.
32
Acciai
INDUSTRIAL-SCALE TESTS
Following the results of the lab tests and in keeping with
specific plant requirements, various tests were performed
in a real plant to evaluate the capability of in-line heat treatment processes (QTB) in the production of high yield
strength rebars.
These tests were essential in order to validate the results
of the laboratory tests, overcome their limitations and simplifications and determine the stability of the in-line process in a real rolling mill.
Based on the results of the laboratory tests, an initial chemical composition was selected in order to ensure good
material weldability.
Before running the tests, the technological parameters
of the rolling mill were studied using the thermal/metallurgical software DLPP (Danieli Long Products Predictor),
which was also used to evaluate the test results [8].
Figure 7 shows the heat profile of one of the tests, from
reheating furnace exit to the cooling bed.
For each test, a suitable number of samples was selected
for technological and metallurgical characterization. The
bend tests and elongation measurements were done using
several methods described in various international standards. Figure 8 shows the effects of bend tests on some
samples, according to various standards.
The rolling tests in conjunction with metallographic and
technological characterization made it possible to deter-
Cooling Time Temperature
[s]
[°C]
1.0
635
2.0
3.0
4.0
526
430
313
Microstrucutural characteristics for composition 2#
Sub-surface
Core
area
Bainite in ferritic
Martensite + PF pearlitic strufture
Martensite + PF
Bainite
Martensite Martensite + PF
Martensite + PF Martensite + PF
Sub-surface
Core
area
Bainite in
Martensite + PF ferritic pearlitic
strufture
[s]
[°C]
1.0
640
2.0
471
Martensite + PF
3.0
4.0
430
277
Martensite Martensite + PF
Bainite and
Ferrite
PF = Proeutectoid ferrite in trace amounts (<5%)
Fig. 4 – Results from microstructural analysis of rebar
core with composition #2 and microhardness profiles
within the cross-sectional area of the quenched rebar,
at increasing cooling times from 1” to 4”.
[s]
[°C]
2.0
558
3.0
4.0
5.0
410
359
288
Sub-surface
Core
area
Martensite +
Bainite
Bainite
Martensite + FP Martensite + PF
within the cross-sectional area of quenched rebar, at
increasing cooling times from 1” to 4”.
within the cross-sectional area of the quenched rebar,
at increasing cooling times from 1” to 5”.
33
Memorie
Fig. 7 – Thermal simulation
using DLPP: temperature
trend of the bar from
reheating furnace exit to
cooling bed entry.
Fig. 8 – Bend tests done on
samples of high-tensile rebar
heat-treated in line.
mine the limitations of the QTB process in the production
of high yield strength rebar, in terms of cooling method,
operating flow rates/pressures, chemical compositions,
productivity and process stability.
It is important to note that for rebars with the same final technological properties, the processing temperatures necessarily
differ depending on the composition. In fact, core hardening is
necessary in some cases in order to reach the desired figures.
This aspect must be taken into consideration to determine
the risk of brittle phases being generated, and the possible
creation of cracks (enhanced by the quenching process). Figure 9 shows a series of macrographs (hardening depth) and
micrographs (surface, core): one of them highlights a crack
generated by a defect that spread within the bar.
Just like in the lab tests, the microhardness profiles in various processing conditions were examined. This made it
possible to determine the exact hardening depth and the
effect of the metallurgical transformations.
CONCLUSIONS
The lab experiments made it possible to assess the behavior of 3 different chemical compositions after subjecting
DIA 16mm rebars to hardening and self-tempering.
34
Fig. 10 – Yield strength measured with increasing
cooling times of rebar in lab tests.
The industrial-scale tests made it possible to evaluate the
performance of the QTB process in the production of high
yield strength rebar (greater than 1000MPa) in terms of
operating flow rates / pressure, optimized chemical compositions, productivity and process stability.
In some DANIELI plants, the QTB process is already being
used for the production of high-strength steels.
Acciai
Fig. 9 – Examples of macrographs and micrographs (surface, core) for high yield strength rebar treated with QTB
at different processing temperatures. Note the sample with a crack generated by a defect.
BIBLIOGRAPHY
[1] M.MIYAJIMA, The Japanese Experience in Design and
Application of Seismic Grade Rebar, Proceedings of Int.
Seminar on production and Application of High Strength Seismic Grade Rebar Containing Vanadium, page 12,
Beijing 2010
[2]S. MORITA, S. HITOSHI, Development of high strength
mild steel deformed bars for high performance reinforced concrete structural members, paper No. 1742. 11th
world conference on Earthquake Engineering, ISBN: 0
08 042822 3, Elsevier 1996
[3]WJE Wiss, Janney, Elstner Associates, Inc., Mechanical
Properties of ASTM A1035 High Strength Steel Bar Reinforcement, Final Report WJE No. 2008.9901.0, 2008
[4]Y. CAIFU, Development of High Strength Construction
Rebars, Proceedings of Int. Seminar on production and
Application of High Strength Seismic Grade Rebar Containing Vanadium, page 58, Beijing 2010
[5]J. HUAIZOHONG, Y. CAIFU, Z. YONGQUAN, Strengthening effects of nitrogen on 20MnSi rebar containing
vanadium, Special Steel, Vol.21,No.5, page 20, 2000
[6]WENG YUQUING, Ultra fine grain steel, Metallurgical
Industry Press, Beijing, 2003
[7]ENGINEERING NEWS RECORD, High strength rebar called revolutionary, Aug/Sept, 2007
[8]C. FABBRO, M. CIMOLINO High Carbon Grades for Wire
Rod Lines - The Core of Danieli Technology, AISTech
2014 Iron & Steel Technology, Indianapolis IN, USA,
May 2014
35
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Metalli non ferrosi
On the ageing of a hyper-eutectic Zn-Al alloy
A. Pola, M. Gelfi, G. M. La Vecchia, L. Montesano
Zinc alloys are widely used in different fields, like handles and locking, fashion and design as well as automotive
or electronics, thanks to their good mechanical and technological properties combined with low cost and easy
formability. A limit to a wider use of these alloys is the aging phenomenon that causes a drop in their mechanical
properties in time. In order to improve their use in competition with more expensive copper and aluminum alloys,
in the last years the research has been addressed to develop new Zn-alloys compositions. One of these new
alloys, containing 15 wt% of Al and 1 wt% of Cu, appears to be suitable for both foundry and plastic deformation
forming processes, as resulted from preliminary laboratory and industrial trials. Being a newly developed alloy,
many properties have still to be investigated, to better understand the effective potentiality for a proper industrial
application. In this paper the ageing behavior of die-cast Zn-15Al-1Cu hyper-eutectic alloy was studied by means
of tensile tests and microstructural analyses. It was demonstrated that the alloy suffers from a drop in mechanical
properties, in particular at the very beginning of soaking at high temperature. A first analysis of the microstructure
by optical and scanning electron microscope was not able to fully point out the causes of the aging phenomenon
Keywords: Zinc alloys - Aging - Die-casting - Mechanical properties
INTRODUCTION
Zinc aluminum alloys are used for the production of several
functional items like hinge and small gears, and also design
parts as handles, taps and fittings. These materials, in fact,
guarantee a good compromise between performances and
production costs. They offer good corrosion and wear resistance. Additionally, zinc alloys can be coated with all the
traditional deposition techniques improving their aspect
according to the specific application [1].
Nowadays zinc aluminum alloys are also good candidates
for the all the application fields where moderate mechanical resistance is required, for example as a substitutive
material for bronze and brass components in home furnishing and fashion parts.
Zn-alloys can be also used for the production of bearing
and bushings as they exhibit for such application high
hardness, wear resistance similar to that of bronze and
many properties comparable to those of cast steel [2-3].
Moreover these materials have a good machinability [3]
that makes easy the manufacture of finished components
with the correct tolerances.
The presence of aluminum, in proper percentage, promotes the fluidity of the alloy, increases the mechanical
properties and enhances the corrosion resistance in mild
aggressive environments. Copper, in percentage between
A. Pola, M. Gelfi, G. M. La Vecchia, L. Montesano
Università degli Studi di Brescia - DIMI
1 and 3 wt%, improves hardness, tensile strength, wear
resistance and creep behavior of the material [4].
The main restrictions in using zinc alloys in structural applications are the high density (more than twice of aluminum and similar to that of steel) and the considerable drop
of mechanical properties with temperature [5-6]. For this
last reason the threshold temperature for the use of zinc
alloys is commonly fixed at 80°C.
A further problem affecting zinc alloys is the formation of
instable phases during solidification and cooling, able to
evolve in time, also at temperature below 100°C [3]. Solid
state transformations through much stable configurations
can occur, causing a decrease of mechanical properties
in time [5]. This drop of zinc alloys properties is normally
known as aging.
Studies on conventional Zn-Al alloys aging showed that
the phenomenon is enhanced by temperature. In particular, mechanical tests performed maintaining sample
at 105°C for different times, showed that the maximum
drop of properties occurs in the first 24h while for longer
soaking times, the decrease is slower and the mechanical properties tend to set at constant values [6]. Based on
these findings some components are artificially aged (24h
at 105°C in the case of automotive parts [5]) before their
use, in order to stabilize the properties and avoid further
drop during the component life.
Fig. 1 shows the Zn-Al binary alloys phase diagram, where the main families of nowadays used Zn-alloys are highlighted. The most common alloys are the so called Zamak
(Zamak2, 3 and 5), characterized by an aluminum content
37
Memorie
of 4% and a copper percentage in-between 0 and 3wt%.
The low melting temperature of Zamak, close to 390°C,
allows the use of hot chamber high pressure die-casting
(HPDC). The hexagonal close packed lattice of these alloys
makes them scarcely deformable and not suitable for hot
stamping [7].
Other commercial alloys with higher aluminum and copper
content are those belonging to the ZA family (ZA8, ZA12,
ZA22, and the most widely used ZA27), which are characterized by higher mechanical properties than Zamak family
[4]. The melting temperature of these alloys is higher than
Zamak ones (over 410°C) and, for this reason, they should
be cast only in cold chamber HPDC or in gravity or sand
casting, with higher cycle time and costs.
Recently, a new zinc alloy with an aluminum content of
15wt% has been developed. This alloy is characterized by
good mechanical properties and high corrosion resistance;
it is also suitable for foundry (hot chamber HPDC) as well
as hot stamping processes [6]. However, many properties
are still not known or under investigation. The aim of this
paper is to study the effect of artificial aging, carried out
at different temperature and for different holding times,
on the tensile properties of this new alloy.
MATERIALS AND METHODS
Tensile test samples in Zn-15Al-1Cu alloy were produced
by hot chamber HPDC with a 500 ton machine. The alloy
was injected into the die at a temperature of 490 ± 10°C,
while mold temperature was set at 250°C. The obtained
samples have a round section with a diameter of 9 mm.
Tensile tests were carried out at room temperature according to UNI EN ISO 6892 standard with a INSTRON 8501
machine, set in displacement control mode, with a cross
head speed of 0.5 mm/min. Displacement was monitored
by an extensometer.
Three samples per condition were tested.
Moreover, on one sample for each aging condition, the elastic modulus was determined with five load-unload cycles
in the elastic field, setting a test speed of 0.1 mm/min.
The evolution of the mechanical properties during aging
was studied at three different temperatures: 80, 105 and
130°C, and for different holding times up to 240h, as reported in Tab. I.
In the time between the alloy production and the aging treatment, all the samples were maintained at a temperature
of -14°C in a freezer, in order to avoid any natural aging.
The “as-cast” samples were tested just after production to
evaluate the mechanical properties in not aged conditions
(0 hours of aging time, as referred in Tab. I).
Three groups of samples were considered, as a function
of the heat treatment temperature. The temperature of
105°C (group A) was selected according to literature findings; in fact, Leis et al. evidenced that, for conventional
zinc alloys, the Ultimate Tensile Stress (UTS) of specimens
aged at this temperature is similar to that of samples naturally aged for one year [4]. Additionally, as already mentio-
38
Fig. 1 - Zn-Al phase diagram [8].
ned, many car producers suggest a pre-treatment of zinc
components at 105° for 24 hours before their use, in order
to stabilize the mechanical properties [5].
The other two temperatures investigated in the present
research were chosen 25°C above and below the reference temperature of 105°C. The temperature of 130°C
(group B) was set to study the drop of properties at higher
temperature, that probably happens with a faster kinetics.
On the other hand, the temperature of 80°C (group C) is
commonly accepted as the upper limit for the use of zinc
alloys [9].
Time [h]
Time [h]
Time [h]
0
0
0
1
2
1
2
8
2
4
24
4
GROUP C:
GROUP B:
GROUP A:
8
168
8
T = 80°C
T = 130°C
T = 105°C
15
240
24
24
168
72
240
168
240
Table. I - scheme of the tensile tests.
It is worth noting that the maximum aging time of 240
hours was chosen to be ten times higher the one considered to be necessary to reach almost stable mechanical
properties for the traditional zinc alloys [6].
In addition to the mechanical characterization, a microstructural analysis was performed on all samples polished
and etched with Nital 2% by means of a Reichert-Jung
MeF3 optical microscope, equipped with the Leica QWin
image analyzer software. The chemical composition of ZnAl phases was assessed by means of the Oxford Energy Dispersive Spectroscopy (EDS) microprobe, coupled to the
LEO EVO 40 Scanning Electron Microscope (SEM).
Metalli non ferrosi
GROUP
t [h]
0
1
2
4
8
15
24
72
168
240
A
T=105°C
261 ± 1,4
248 ± 1,0
248 ± 1,0
245 ± 3,5
246 ± 6,4
236 ± 0,1
238 ± 4,6
228 ± 4,7
224 ±3,6
YS [MPa]
B
T=130°C
271±6.4
242 ± 1,5
239 ± 3,6
225 ± 0,7
210 ± 3,5
196 ± 2,1
192 ± 1,2
C
T=80°C
A
T=105°C
251 ± 4,2
255 ± 1,4
246 ± 0,7
246 ± 0,8
300 ± 7,2
310 ± 6,6
310 ± 6,6
304 ± 0,4
296 ± 3,3
295 ± 5,5
290 ± 3,8
281 ± 2,5
279 ± 3,5
290 ± 4,5
249 ±5,1
246 ± 1,4
UTS [MPa]
B
T=130°C
313±7.4
306 ± 6,7
299 ± 3,1
281 ± 1,3
271 ± 2,4
252 ± 3,6
248 ± 0,4
C
T=80°C
300 ± 0,9
312 ± 4,6
304 ± 0,4
295 ± 6,0
290 ± 2,8
298 ± 4,8
301 ± 6,3
Table. II - tensile tests results.
RESULTS
TENSILE TESTS
Table II shows the values of the UTS and Yield Strength
(YS) of Zn-Al alloy samples artificially aged at different times and temperatures.
Fig. 2 shows the behavior of UTS as a function of the aging
time. Similarly to what reported in literature for other zinc
alloys [3], the new composition suffers from a strong drop
of the mechanical properties in the first 24 hours of aging.
Results from samples aged for longer time, however, reveal a further, even though slower, reduction of the UTS. The
behavior of the YS is similar to that of the UTS.
For the same aging times, group B samples, i.e. those treated at higher temperature, show lower UTS than group
A samples and the reduction of the mechanical properties occurs faster. Comparing the drop of the UTS, group
A samples has a drop of 6% in the first 24 hours and 11%
after 240 hours. For the group B samples the reduction is
10% in 24 hours and 21% after 240 hours. Finally, in the
case of samples C, aged at 80°C, the maximum drop of
the UTS after 240 hours is only 4%.
As shown in Fig. 2, the aging at 105°C for 240 hours gives
similar UTS and YS values to those obtained at 130°C for
24 hours. From this experimental result, it follows that the
same mechanical properties can be achieved lowering the
aging time and increasing conveniently the heat treatment
temperature. Such aging conditions can be reasonably accepted only if the cast part is not affected by high levels of
porosity or other defects that can be modified or worsened
by a long stay at high temperature. It should be also taken
into account that a soaking at high temperature usually causes a a coarsening and globularizarion of the precipitates.
Finally, a last aging test was performed at 105°C for 720
hours (30 days), aimed at evaluating if a steady state condition of mechanical properties can be finally reached. The
UTS and YS measured after this heat treatment were 245
MPa and 193 MPa respectively, which are values close to
those obtained after the aging at 130°C for 240 hours,
confirming the achievement of almost stable conditions.
Fig. 2 - Ultimate tensile strength of Zn-Al samples
versus aging time.
A conclusion that can be derived from this result is that a
heat treatment at 105°C for 24 hours is not enough to reach
a complete settlement of mechanical properties for a Zn-15Al-1Cu alloy. To foresee the behavior of a component designed with this alloy for long service applications, the aging
treatment should be performed at higher temperatures.
Additionally, tests carried out on samples aged at 80°C
proved that the drop of mechanical properties is strongly
slowed down by the reduction of the treatment temperature, confirming that 80°C can be assumed as the threshold
temperature also for this alloy
MICROSTRUCTURE
In order to better understand the aging phenomenon, a
metallographic characterization of as-cast and aged samples was performed by means of optical microscope. Ac39
Memorie
Fig. 3 - Microstructure of Zn-15Al-1Cu alloy die-cast samples a) “as-cast”, b) after 240 h of aging at 105°C and c)
after 240 h of aging at 130°C.
cording to the phase diagram (Fig. 1), the microstructure
of the Zn-15Al-1Cu alloy should be composed at room
temperature by Al-rich primary phase and Zn-rich eutectic,
both transformed by the eutectoid decomposition.
Figure 3 shows the microstructures of samples in different
conditions: as-cast (a), aged for 240 hours at 105°C (b)
and at 130°C (c). The Al-rich primary grains (dark areas)
are surrounded by the eutectic phase (light gray areas). No
particular differences can be noted between samples, in
terms of size, shape or amount of the different phases. No
recrystallization phenomena are detectable.
All samples are almost free from porosity, regardless of
the aging conditions, which means that the reduction of
the mechanical properties cannot be ascribed to an increase of defects, as can easily occur in die-cast components
after long stays at high temperature.
Preliminary SEM analyses were carried out on the samples
in the as cast condition (Fig. 4a) and heat treated for 240
h at 130°C (Fig. 4b). Both samples show the eutectoid decomposition of Al-rich primary grains with the consequent
formation of a thin lamellar microstructure, surrounded by
the Zn-rich eutectic. The Al content of primary dendrites is
close to the theoretical value of 32.3% expected from Zn-Al
phase diagram. This suggests that notwithstanding the ra40
pid cooling imposed by the production process (HPDC), the
microstructure is not far away from the equilibrium state.
Comparing figure 4a and figure 4b some small differences
can be appreciated in the microstructure. In particular, the
aged sample shows more globular and coarser precipitates at the interface between the phases and the eutectoid
lamellae seems to be better defined, maybe as a consequence of a more distinct separation between Al and Zn
elements.
Such rearrangement of the microstructure could justify the
drop of mechanical properties measured by tensile tests
on aged samples. Increasing the temperature, in fact, the
atomic mobility is enhanced and solid state diffusion and
solute redistribution are promoted. The achievement of a
pseudo-equilibrium state at the end of the aging treatment
should allow the formation of a structure free from supersaturated solutions, having a lower local lattice deformation. This condition causes a reduction of barriers to the
dislocation movement and a modification of the free mean
path and, consequently, a lower mechanical resistance.
Unfortunately, such modifications are not detectable by
optical or electron microscope and also the spatial resolution of EDS microanalysis is inadequate to measure a
solute redistribution on so small distances. Transmission
Metalli non ferrosi
Spectrum
Al
Cu
Zn
Spectrum
Al
Cu
Zn
1
32.03
0.39
67.58
1
36.48
0.17
63.34
97.07
2
8.98
0.55
90.47
2
2.11
0.82
Fig. 4 - SEM analysis performed on the as cast sample (a) and sample treated for 240 h at 130°C (b).
Electron Microscope (TEM) measurements and X-ray diffraction (XRD) experiments are desirable to overcome
such limitations and explain the results of tensile tests in
terms of the alloy microstructural changes.
le tests, anyway. To examine more deeply this hypothesis
further investigations with other techniques like TEM and
XRD are under development.
ACKNOWLEDGMENT
CONCLUSIONS
In this paper the aging behavior of the hypereutectic Zn15Al-1Cu alloy was investigated.
Being a newly developed alloy, no data are up to now available in literature about natural or artificial aging. Tensile
tests were performed on samples aged at three temperatures (105, 130 and 80°C) for different soaking times (up
to 240 hours and in one case up to 720 hours) to measure
the decrease of mechanical properties.
This characterization gave useful information about the
potential use of this alloy at relatively high temperatures
and allowed to assess the mechanical properties close to
steady state conditions.
The results show that in the first 24 hours a fast decrease of mechanical properties occurs at all the investigated
temperatures. After this period the reduction of mechanical strength slows. As expected, the aging phenomenon
is accelerated by increasing the aging temperature: the
reduction of the UTS after 240 hours was 4% at 80°C, 11%
at 105°C and 21% at 130°C.
Microstructural analyses carried out by optical microscope and SEM-EDS revealed only small differences between
the aged and the as-cast samples. In particular, the aged
samples showed a slight coarsening of precipitates at
the phases boundary and the eutectoid lamellae appeared better defined, suggesting the redistribution of solute
elements. Such microstructural modifications could justify
the reduction of mechanical properties measured by tensiLa Metallurgia Italiana - n. 4/2015
The authors want to thank Mrs. Valentina Ferrari for the
support in the execution of tensile tests.
REFERENCES
1. D. APELIAN, M. C. Donald HERRSCHAFT, Casting with
zinc alloys, JOM 33 (11), 1981, 12-20.
2. K.J. ALTORFER, Zinc-alloys compete with bronze in bearings and bushings, Metal Progress 122 (6), 1982, 29-31.
3. E. J. KUBEL, Expanding horizons for ZA alloys, Adv. Mat.
Proc. inc. Material Progress, 1987, 51-57.
4. Y. H. ZHU, W. B. LEE, S. TO, Ageing characteristics of
cast Zn-Al based alloy (ZnAl7Cu3), J. Mat. Sc. 38 (9),
2003, 1945-1952.
5. A. BUCCIOL, Aging behaviour of zinc die casting alloy
ZP0810 (Master thesis), Univesity of Padova, 2012.
6. W. LEIS, L. KALLIEN, Ageing and creep of Zinc-Diecast
alloys, Int. Zinc Diecasting Conference 2013 “Tradition
& Innovation”, Praha 13-14 June 2013.
7. A. POLA, L. MONTESANO, R. ROBERTI, Nuove leghe di
zinco per l’industria del design, Proceedings 33° Convegno Nazionale AIM, Brescia 10-12 November 2010.
8. Alloy Phase Diagrams, ASM Metals HandBook Vol. 3,
ASM International 10 ed., 1992.
9. A. POLA, R. ROBERTI, D. ROLLEZ, Primary and steady
state creep deformation in Zamak5 die-casting alloy at
80°C, Mat. Charact. 59 (12), 2008, 1747-1752.
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Forgiatura
Implementation of an open-die forging
process for large hollow shafts for wind
power plants with respect to an optimized
microstructure
M. Wolfgarten, D. Rosenstock, L. Schaeffer, G. Hirt
To realize large wind power plants in an economically feasible way, it is necessary to identify potential for
lightweight design of the generator hollow-shafts, which are commonly produced by casting. The weight of
these shafts can significantly be reduced by producing them by open-die forging, since the forming of the
material leads to a higher strength, which allows to reduce the wall thickness noticeable. This paper describes
the development and implementation of a forging process for hollow shafts with respect to an optimized
microstructure. To numerically investigate this process, a realistic finite element simulation model was
developed in a first step. The kinematic of the tools has been implemented authentically to provide a realistic
material flow and process conditions. Additionally, a material model for the steel 42CrMo4 was integrated into
the simulation model to predict the resulting microstructure. Using the implemented FE model, the forging
process was optimized manually to achieve a homogeneous and fine-grained microstructure. The optimization
was based upon a variation of different forging parameters and the sequence of forging steps. In the next
step, a forging on laboratorial scale was performed to validate the simulation model. For this purpose, after
forging, specimens from the hollow shaft were evaluated by metallography to determine the final grain size. A
comparison of the results with the numerical simulation showed a general agreement of the measured grain
size with the numerically calculated grain size. Based upon these results, the process model was transferred
to an FE model with an industrial scale. By this it was possible to analyze the transferability of the used FE
model regarding the assumptions about the kinematics and the sequence of the forging steps. A numerical
investigation of the industrial process proved the scalability of the process to an industrial relevant geometry.
Keywords: Open-die forging - Microstructure - Process optimization
Introduction
Initial Situation
During the last years the importance of alternative energy
sources like wind energy has risen constantly. To meet
with the further anticipated demand for renewable energy,
the importance of wind energy will be growing even
more as this currently is the cheapest and most effective
renewable energy source [1]. A higher energy production
by wind power plants cannot only be achieved by setting
M. Wolfgarten, D. Rosenstock, G. Hirt
Institute of Metal Forming,
RWTH Aachen University
L. Schaeffer
Laboratório de Transformação Mecânica, Universidade
Federal do Rio Grande do Sul, Brazil
up additional wind parks, but also by an increase in the
performance and by this the size of newly built plants.
This requires larger machine parts, like rotors, shafts,
gears and generators, leading to higher requirements for
the tower’s construction. To cope with these aspects, the
reduction of the nacelle’s weight at the top of the tower
offers a good opportunity. One approach is to replace the
commonly cast generator shaft by a forged hollow shaft with
excellent mechanical properties. In comparison to a cast
hollow shaft, a forged shaft could offer a higher strength
and therefore would allow reducing the wall thickness
significantly. Recker et al. [2] estimated that producing
the hollow shaft by open-die forging could allow a weight
reduction of up to 60% compared to a cast hollow shaft.
State of the Art
Open-die forging is mainly used for the production of high
quality parts in low quantities. These parts are applied for
43
Memorie
highly loaded purposes in large machines like generator
shafts or rolls. Hollow parts like rings are usually forged
by applying an upper-die and a mandrel. Usually, this
kinematic is used for forging rings and intends to increase
the diameter. For the production of a hollow shaft, which
has a much higher length and smaller diameter, the
main challenge is to ensure a sufficient axial material
stretching. The longitudinal material flow is predominantly
achieved by improving the tool’s shape and the process
kinematics [3]. The best axial stretching can be realized
using concave or v-shaped dies. However, two concave
dies are lacking flexibility during the process according
to the forgeable diameter as the shape of the dies limits
the possible workpiece geometries. Two v-shaped dies
are disadvantageous for the process handling, since this
tool combination impedes the loosening and extracting of
the mandrel after a forging heat. Based upon numerical
investigations and literature review [4], the combination of
a flat and a v-shaped die proved best to realize a sufficient
axial stretching.
Motivation and Scope
The overall objective of the investigation presented in
this paper is the optimization of the open-die forging
process for hollow shafts with respect to an optimized
microstructure and transferring the results to an
industrially relevant geometry.
As a first step, the main focus was the development of a
simulation model for the forging of a 150 kg workpiece
under consideration of the kinematic and tool geometry.
The second main objective consisted of optimizations
of the microstructural properties meant to ensure the
final product’s high mechanical strength. This requires
a homogeneously distributed, fine grain size along the
workpiece, which is likewise supported by a homogeneous
strain. So a manual optimization through variation of
the bite ratio and height reduction was performed.
To verify the results from the numerical simulation,
especially in terms of microstructural evolution, forging
of a 150 kg workpiece was performed at the IBF. The
corresponding results from the experiment were used to
validate the numerical simulation model according to the
microstructural evolution, tool geometry and kinematics.
Based upon the validated simulation model, in a third step,
the process was transferred to an industrially relevant
geometry by numerical investigation.
Methods and Procedure
Numerical Simulation Model
For correctly predicting the material flow and the
microstructural evolution, a simulation model was
implemented in Transvalor Forge 2011, whose boundary
conditions and kinematics coincide with the forging
conditions in reality. The general requirements on the
process kinematic for a realistic simulation model are:
1. The upper flat die is moving in y-direction and
44
performs the reduction of the material.
2. The lower v-shaped die is fixed in its position and not
moved during the whole process.
3. The mandrel supports the hollow shaft during the
process and is held by a manipulator. When the upper
die presses in y-direction, the flexibly supported
mandrel can move freely in all directions.
The grain size calculation was realized by implementing
a microstructure model into the integrated calculation in
Forge 2011, based on experiments at IBF and literature
[5-8], which showed good accordance with the behavior
of the used material.
Fig. 1 - Process setup – hollow shaft forging
The numerical simulation of hollow shaft forging is more
complex compared to conventional open-die forging due
to positioning of the workpiece over a mandrel and the
attachment of the mandrel to the manipulator. This setup
and the long process time increases the complexity
and calculation time of the simulation significantly. The
simulation of the reference process described within the
paper requires a calculation time of one week on a quad
core Intel Xeon workstation. Furthermore, the handling of
the workpiece in the simulation is impeded since in hollow
shaft forging the workpiece is just indirectly positioned
during the process. Sliding of the hollow shaft leads to an
inexact positioning. Hence it is not possible to simulate the
whole forging process in one simulation. The simulations
needs to be interrupted after every pass to control the
exact positioning and kinematic.
Chosen Strategy for Hollow Shaft Forging
Different parameters influence the forging of a hollow
shaft. Firstly, the main influencing parameters are standard
parameters for open-die forging, the height reduction εh
and the bite ratio sB/h0, which describe the ratio of the
contact length between die and workpiece and the initial
height (here: initial diameter). The variable parameters are
the direction of forging and the combination of forging
strokes and rotations. According to the forging direction,
two different possibilities exist. At first, the workpiece
can be forged from the front of the mandrel towards the
manipulator tong. After one rotation has been forged,
the manipulator feed is executed and the next rotation is
forged.
The second possibility consists of forging the opposite
direction. As described before, after the forging of one
Forgiatura
rotation, the manipulator feed is executed and the next
rotation is forged.
Besides the forging direction, the combination of rotation
and strokes can be varied as second possibility to mainly
influence the forging process. The investigations described
within this paper are based on the strategy to forge one
rotation at first, translate the workpiece by the manipulator
feed and forge the second rotation in the following step.
Figure 2 visualizes this process principle. As shown in
the top, the whole circumference of the hollow shaft is
forged in the first step, which requires 10 strokes or a
whole workpiece rotation. After each stroke, the mandrel
is rotated by 90°, but just after the fifth stroke of each
rotation the mandrel is once rotated by 45°.
Process optimization
Generally, the bite ratio (quotient of the bite length and
initial height of the workpiece – sB/h0) and the height
reduction εh can be identified as the forging parameters,
which most decisively influence the strain distribution
and grain size. Therefore to optimize the process
parameters for hollow shaft forging, numerical studies of
different bite ratios (0.3, 0.5, 0.7) and height reductions
(10%, 20%) were performed and the resulting equivalent
plastic strain and average austenitic grain size were analyzed.
The optimization was performed for a hollow cylinder with
a diameter of 240 mm at an initial temperature of 1200 °C.
For the forging on a laboratorial scale, this corresponds to
the geometry for the middle steps of the hollow shaft.
As due to the long simulation time the process could not
be optimized completely, three different process routes
were used to investigate the process optimization. The
optimization was performed for the forging of two rotations,
each for one and two passes. The equivalent strain and
the grain size according to the optimization are evaluated
along three lines at half of the wall thickness, each 120°
distributed over the circumference, see Figure 3.
Table 1 summarizes the influence of the bite ratio, height
reduction and number of forging steps on the strain
distribution in the workpiece. An increase of the bite ratio
from 0.3 to 0.7 leads to an average increase of strain of
26% for one pass and 50% for two passes at half of the wall
thickness. The standard deviation of the strain distribution
can be reduced by 46% for one and by 14% for two passes
and thus mainly increases the strain homogeneity. Similar
effects can be observed for a higher height reduction.
εh/Number of
Passes
Parameter
Fig. 2 - Strategy for the forging of one rotation [2]
Fig. 3 - Points for evaluation of strain and grain size
distribution
So for the optimization of the strain, a higher bite ratio
and height reduction should be preferred to achieve the
intended results.
The grain size as second optimization objective is
mainly influenced by the temperature, strain and strain
rate. Since the minimum warm forming temperature for
42CrMo4 is 850°C, the workpiece temperature should
not drop below this point. So to allow a long enough time
frame for forging the initial temperature is set to 1200°C.
Bite ratio (sB/h0)
0.3
0.5
0.7
10% / one pass
0.46 ± 0.13
0.56 ± 0.13
0.58 ± 0.07
20% / one pass
0.54 ± 0.11
0.97 ± 0.14
1.03 ± 0.13
10% / two passes
0.75 ± 0.25
1.03 ± 0.14
1.13 ± 0.14
Table 1 - Influence of bite ratio, height reduction and forging steps on the equivalent plastic strain distribution
(average and standard deviation)
45
Memorie
Properties
optimization
Bite ratio
Parameter
0.3
0.5
0.7
10% / one pass
86 ± 21 µm
56 ± 21 µm
55 ± 20 µm
20% / one pass
82 ± 31 µm
49 ± 16 µm
40 ± 16 µm
10% / two passes
63 ± 17 µm
47 ± 17 µm
30 ± 6 µm
Table 2 - Influence of parameters on average austenitic grain size (average and standard deviation), initial grain
size 200 µm
Therefore, a variation of the temperature to optimize the
microstructure is nearly impossible as always the maximum
temperature of 1200°C has to be chosen for the beginning
of the forging process. The strain rate results from the tool
speed of 40 mm/s and the geometry of the work piece.
Table 2 visualizes the influence of the forging parameters
on the average austenitic grain size evolution during the
forging process. It can be concluded that a large bite ratio
in combination with a high height reduction is preferred to
yield a fine grained and homogeneous microstructure.
From this case study, a process route with a higher bite ratio
and a sufficiently height reduction should be preferred to
both optimize the strain distribution and microstructure in
the hollow shaft. Furthermore, the numerical simulations
proved that a height reduction of 20% and more leads to
an increase of the inner diameter and deviations from the
intended shape. Therefore, the height reduction during the
process should be lower than a value of 20%.
Besides a small and homogeneous microstructure along
the workpiece, likewise a homogeneous distribution over
the cross-section is preferred. As example, Figure 4 shows
the grain size distribution after the forging of one rotation.
It can be concluded that the chosen forging strategy is
advantageous to achieve a homogeneous strain distribution
over the cross-section. The strong deviation in two points
is probably caused by numeric irregularities and disagrees
with the general distribution at the cross section.
Fig. 4 - Exemplary grain size distribution at the cross
section
Reference forging process
The reference forging process was used to validate
the numerical simulation model and in particular the
microstructure calculation. Therefore the process
was designed such that an inhomogeneous grain size
distribution should be obtained in the final workpiece to
validate the microstructure calculation on a preferably
wide range. For that purpose, sections with different
outer diameters were forged using different reductions in
diameter and varying reheating conditions. The geometry
for the laboratorial forging was limited by the maximum
possible length of 490 mm and a weight of 150 kg for the
initial workpiece. The development of the shape of the
Fig. 5 - Geometry in different
stages of the reference process,
all dimensions in mm
46
Forgiatura
workpiece during forging is shown in Figure 5.
The forging was performed using a 6.3 MN hydraulic press
and a 6-axis forging robot. The workpiece was initially
heated to 1200 °C and the tool speed was set to 40
mm/s.
Results
Results of the forging process
Comparing the final shape, material flow and especially the
grain size of the experiment to the numerical simulation
allows judging the quality of the simulation model. In
this context, the validation of a simulated grain size by
metallographic analysis is limited since only the final
state can be investigated. So it is not possible to analyze
the grain size evolution during intermediate steps of the
forging process.
Figure 6 visualizes the results of the metallographic
analysis. Generally, the microstructure is only investigated
in one single center point per step and could possibly vary
due to different conditions in a step. The metallographic
analysis shows that in the largest step (I) the microstructure
is just partially recrystallized and shows still similarities
to the initial microstructure. This effect results from the
low strain, which just has been imposed by the 10 mm
diameter reduction during the first pass. In the other three
grain size and recrystallized microstructure, is fulfilled.
According to the numerical calculation, during the
reheating process, the grain size increases rapidly up to
approx. 500 µm due to grain growth, whereas the dynamic
and static recrystallization phenomena during and after
a strain increment result in a reduction of grain size.
Considering the standard deviation of the metallographic
grain size measurements, the grain sizes show roughly the
same tendency as summarized in Table 3.
A correlation between the strain and the grain size can
generally be observed, showing that a higher strain
leads to a finer grain size in the workpiece. However, this
needs always to be regarded in relation to the process
history, since significant grain growth can appear during
a necessary reheating process. For this kind of process,
an exact prediction of microstructure proves as difficult
due to the long process time and the semi-empirical
modelling of the microstructure evolution. Nevertheless,
the microstructure calculation reproduces the influences
on microstructure during forging in a satisfying way.
Transfer to Industrial Scale
The verification of the simulation model showed in
principle a satisfying accordance between the real process
and the numerical model. Based upon this, the simulation
Fig. 7 - Industrial hollow shaft geometry, all
dimensions in mm
Fig. 6 - Exemplary results of microstructure
investigation for each step
steps of the shaft, the effect of the higher strain clearly
becomes visible as the grain size becomes smaller in a
fully recrystallized microstructure. Under consideration of
the measuring error, the average austenitic grain size for
each step is decreasing from 142 µm in step I to 37 µm for
step IV with the smallest diameter. The general expectation
from the optimization, that a higher strain leads to a finer
Step
1*
2
3
4
Metallography 142 ± 89 µm 102 ± 21 µm 110 ± 38 µm 37 ± 14 µm
Simulation
167 µm
163 µm
84 µm
6 µm
Absolute Deviation
17,6%
59,8%
23,6%
83%
in Percent
* The high standard deviation is resulting from the occurrence of non-recrystallized grains.
is scaled to a 20 ton geometry (Figure 7), based on [9],
in order to prove the process design on a large scale. A
general difference between forging on laboratorial and
industrial scale consists in the thermal development
during the process. Due to the significantly larger volume
and a smaller ratio of surface to volume, the workpiece
cools down considerably slower. Therefore, the number of
necessary heats can be reduced from four heats for the
150 kg part to just two heats for the 20 t part.
Fig. 8 - Strain evolution for a hollow shaft with an
industrial geometry
Table 3 - Comparison of numerical and real grain size
47
Memorie
Fig. 9 - Grain size evolution for a hollow shaft with an industrial geometry (Reheating after 3rd pass)
According to the strain distribution in a longitudinal
section of the workpiece, as shown in Figure 8, the
chosen forging strategy leads to a homogeneously and
sufficient equivalent strain in the three smaller steps of
the workpiece with values of εV>5 in average. Only in the
first step (largest diameter), the average strain reaches a
level of εV just above 2. Nevertheless, a homogeneously
distributed average equivalent strain of 2 is sufficient to
enable good mechanical properties.
Figure 9 visualizes the grain size distribution in a longitudinal
section. The 2nd, 3rd and 4th step of the workpiece are
showing a fine and homogeneously distributed grain size
between 5 µm and 40 µm. In contrast, the largest step
of the hollow shaft has a final grain size up to 1200 µm.
The reason for this behavior is based on the chosen
process sequence. The workpiece needs to be reheated
to forging temperature of 1200 °C, after the temperature
has dropped below the minimum forging temperature
of 850°C. As the step I has already been forged in the
first heat and is not forged any more after reheating, the
resulting microstructure is mainly driven by grain growth,
which leads to an enormously high grain size of up to 1200
µm. In contradiction the grain size of step II, III and IV is
mainly influenced by the forging during the last heat. As
shown in Figure 9, the forging on industrial scale results
apart from step I in a more homogeneous grain size than
the forging on laboratorial scale.
Conclusion
This paper proves by numerical simulation and experiments
that open-die forging can generally be used for the process
design of hollow shaft forging. The experimental forging
of a 150 kg part based on numerical simulation studies
showed that the developed kinematic and pass schedule
are suitable.
The development of the experimental forging process was
based on an optimization by variation in the numerical
simulation. The optimization showed that high bite ratios
of 0.7 and height reductions of 10-20% should be preferred
48
for an optimized distribution of strain and grain size in the
workpiece.
In order to validate the numerical simulation model, a 150
kg hollow shaft was forged and compared to an identically
simulated numerical process. A comparison of the
numerically predicted and the experimentally measured
grain size showed a general accordance between both
approaches. However, due to the long process time and the
semi-empiric microstructure modelling in the numerical
simulation, some uncertainties occurred, leading to an
average deviation of 46% between the numerical and
experimental results. Since in general the trend of the
numerical simulation could be observed in the experiment,
the results can be considered in principle as sufficient.
Based on this, the process design of the 150 kg shaft
could be transferred to an industrial scale. A numerical
investigation of the industrial process proved the scalability
of the process to an industrial relevant geometry. Both in
laboratorial and industrial scale, the grain size evolution
is mainly influenced by grain growth. In the industrial
scale, the resulting grain size distribution is much more
homogeneous as less reheating steps are required and
by that, the influence of grain growth on the grain size is
reduced.
Outlook
•The manual optimization presented in this paper only
considers the optimization of few rotations and passes.
To perform a more complex optimization, it is necessary
to regard the whole process. Possible approaches are
a fully automatic optimization through the variation of
different process parameters or a pass-wise optimization
based upon the main values for the investigation of the
microstructure as strain, strain rate and temperature.
As a fully automatic process optimization in numerical
simulation is too complex regarding simulation time, a
fast empirical calculation model could be implemented
and used to optimize whole processes. First works to
this direction for squared ingots are presented in [10].
Forgiatura
•A further degree of freedom to influence the workpieces
final properties is the initial geometry. By a variation
of the height to diameter ratio, the geometrical and
microstructural evolution are influenced significantly.
•The investigation of the grain size evolution showed
that grain growth during reheating has the dominant
influence on the final grain size. Therefore an optimization
approach could be the adaption of the process route, so
that in the final heat the whole workpiece is forged in
order to reduce the large grain size after grain growth.
•While lowering the reheating temperature is not an
option for the 150 kg workpiece, this might be a solution
for the 20 ton ingot. As the numerical simulation is
validated, it could be used to find the optimal reheating
temperature.
Acknowledgements
The authors thank the Deutsche Forschungsgemeinschaft
and CAPES for the financial support within the project
“Forged hollow shafts for power plants” in the “Brazilian
German Collaborative Research Initiative in Manufacturing
Technology” (BRAGECRIM).
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in a low alloy steel“ in “Journal of Material Science” Edition 44, Page 835-842, 2009
7] Y.C. Lin, M.Cheng, “Numerical simulation and
experimental verification of microstructure evolution in
a three-dimensional hot upsetting process” in “Journal
of Materials Processing Technology” Edition 209, Page
4578-4583, 2009
8] D. Qian, J.Guo, Y. Pan, “Austenite Grain Growth Behavior
of AISI 4140 Alloy Steel”, Wuhan, 2010
9] T. Noack and S. Nelle:, Titel: Fehlerfreier Riese: Die
Rotorhohlwelle einer 5 MW-Windkraftanlage. in:
Giesserei, 2008. 95(9): p. 30-34. 2008
10]D. Rosenstock, D. Recker, M. Franzke, G. Hirt, D.
Sommler, K.-J. Steingießer, A. Tewes, R. Rech, B.
Gehrmann, S. Kirchhof, R.Lamm, “Online-Visualisation
during Open Die Forging and Optimisation of Pass
Schedules” in “Steel Research International”, 2013
49
Atti e notizie
Corso itinerante
SOMMARIO
SOLIDIFICAZIONE E
COLATA CONTINUA
VITA ASSOCIATIVA
7-8-14-15-22 maggio 2015
Organizzato da:
Centro di Studio Acciaieria,
ASSOCIAZIONE ITALIANA
DI METALLURGIA
L’Associazione Italiana di Metallurgia ha deciso, come da tradizione,
di organizzare la nuova edizione del
Corso sulla colata continua degli
acciai per continuare a sostenere le
imprese nell’azione di formazione del
proprio personale, uno dei fattori per
mantenere e migliorare la propria
competitività. Data l’importanza dei
semilavorati destinati alle operazioni
di forgia, durante il Corso si affronteranno anche le problematiche relative al colaggio dei lingotti e dei blumi di grande dimensione. Secondo
la formula ampiamente collaudata e
gradita ai partecipanti, la formazione
viene svolta con due modalità: lezioni di tipo teorico - volte a fornire
i concetti di base relativi agli aspetti
metallurgici e al funzionamento degli
impianti - e visite tecniche presso gli
impianti produttivi - in modo da poter
osservare sul campo gli aspetti più
significativi e peculiari riguardanti i
sistemi per la colata continua.
Per rispondere a queste esigenze il
Corso è itinerante e le lezioni si svolgeranno presso alcune interessanti
realtà produttive: Acciaieria Arvedi,
Prosimet, Acciaieria di Calvisano, Duferdofin Travi e Profilati di Pallanze e
Calderys.
Dopo una sintetica introduzione di
carattere storico il Corso abbraccerà
i temi della solidificazione, le problematiche relative alla struttura della
macchina, il colaggio dei lingotti, i
componenti refrattari, le polveri di copertura, l’applicazione dei campi elettromagnetici, la difettologia, i modelli
di simulazione, il colaggio di billette,
blumi e bramme ecc..., solo per citare
i temi salienti. In occasione delle visite i tecnici delle società ospitanti presenteranno gli impianti con una particolare sottolineatura degli aspetti
caratteristici di ogni sistema di colata.
50
Solidificazione e colata continua... 51
Scuola Metallurgia delle polveri...52
Giornata di studio
“La Metallografia passando dalla
preparativa metallografica”.......53
Failures nei refrattari................53
Prossime manifestazioni AIM....54
Queste attività saranno affiancate ed
integrate da interventi didattici tenuti
da docenti del Politecnico di Milano,
nonché da esperti di società di ingegneria di riconosciuta esperienza.
Gli interventi didattici e le visite tecniche si articolano su un arco di cinque
giorni e sono organizzati con cadenza
tale da evitare ai partecipanti un’assenza eccessivamente prolungata
dalle proprie aziende.
Il consistente numero di visite tecniche, che sono state organizzate per
migliorare la qualità del percorso culturale, costringono a limitare il numero dei partecipanti, per cui si consiglia di provvedere all’iscrizione il più
presto possibile.
Il Corso, coordinato da Ottavio Lecis, Silvia Barella e Serena Fasolini,
si svolgerà secondo il seguente programma:
giovedì 7 maggio (Cremona)
• Gli aspetti principali del processo
di solidificazione: strutture di solidificazione e segregazioni
• Modalità di colaggio
• Visita agli impianti di Acciaieria Arvedi
venerdì 8 maggio (Bottanuco BG)
• Polveri di colata continua: che cosa
sono, a cosa servono e come funzionano
Attività dei Comitati
Tecnici...................................55
Notizie da Unisider........ 57
• Refrattari nei sistemi di colata continua
• Magnesita Navarras, Zubiri, Spagna
• Refrattari per la colata in sorgente
• Presentazione di Prosimet
• Visita agli impianti Prosimet
giovedì 14 maggio (Calvisano BS)
• Applicazioni di sistemi elettromagnetici
• Ergolines Lab, Trieste
• Controllo di parametri di processo
in colata continua
• Presentazione di Acciaierie di Calvisano (Feralpi Group)
• Visita agli impianti di Acciaierie di
Calvisano (Feralpi Group)
venerdì 15 maggio (Pallanzeno VB)
• Il colaggio in blummi di gradi dimensioni
• Principi base per il colaggio dei lingotti
• Presentazione di Travi e Profilati di
Pallanzeno (Duferdofin Nucor)
Vita associativa
• Visita agli impianti di Travi e Profilati di Pallanzeno
• Sinterizzazione e sinterotempra
• Nuovi prodotti e futuri sviluppi in
campo automobilistico
• La precisione dimensionale dei particolari sinterizzati
• Requisiti geometrici dei prodotti
sinterizzati multilivello in accordo
con le nuove tecnologie del processo di calibratura
• I materiali sinterizzati: influenza
della porosità sulle caratteristiche
• Le operazioni post-sinterizzazione
• I cuscinetti autolubrificanti
• Visita degli impianti SACMI
• La normativa sui materiali sinterizzati e gli acciai sinterizzati
• Trattamenti termici e termochimici
dei sinterizzati ferrosi
venerdì 22 maggio (Fiorano Modenese MO)
• I difetti del processo colata continua
• Soluzioni impiantistiche per eliminare i difetti per i tipi di acciaio
• Danieli & C. Officine Meccaniche,
Buttrio
Scuola
Metallurgia
delle Polveri
Imola, 9-10 giugno 2015
Nei principali paesi industrializzati la
metallurgia delle polveri è ormai una
tecnologia consolidata. Questa posizione riconosciuta favorisce l’organizzazione di corsi di formazione tecnico-scientifica sui processi e sui requisiti richiesti per una corretta scelta e
una valida progettazione dei materiali
e dei componenti sinterizzati.
Le grandi organizzazioni internazionali di categoria (APMA - Asian Powder
Metallurgy Association, APMI - American Powder Metallurgy Association,
MPIF - Metal Powder Industries Federation, EPMA - European Powder Metallurgy Association) sono i maggiori
enti che, per statuto, sono impegnati
in queste attività, finalizzate a favorire la diffusione delle conoscenze specifiche e gli incrementi delle possibili
utilizzazioni.
In questo panorama, l’AIM, per quasi trent’anni, dal 1960, organizzò dei
corsi di formazione sulle possibilità
e sugli impieghi della metallurgia
delle polveri. Poi, una decina d’anni
prima della fine del secolo, per motivi sostanzialmente legati a forme di
concorrenza esasperata fra gruppi di
aziende del settore, quella bella tradizione fu interrotta.
Grazie all’interesse e alla disponibilità
dimostrati dalla SACMI, è stato possibile ripartire con questa bella iniziativa. Dopo il successo di una prima edizione, nel 2012, l’Azienda di Imola ha
proposto all’AIM di ripetere l’evento,
allo scopo di favorire la diffusione di
10 GIUGNO 2015
conoscenze rigorose e aggiornate su
una tecnologia competitiva, la cui affermazione può contribuire “nel suo
piccolo” alla ripresa dello sviluppo
economico nazionale.
Il Centro Metallurgia delle Polveri
dell’AIM e la SACMI si augurano che
l’impegno profuso per l’organizzazione del Corso – anche da parte dei docenti, tutti specialisti del settore – sia
adeguatamente riconosciuto attraverso un’ampia partecipazione all’iniziativa di tecnici e studiosi interessati
alla metallurgia delle polveri.
La scuola si svolgerà secondo il seguente programma:
9 GIUGNO 2015
• Presentazione della SACMI
• Presentazione del Centro Metallurgia Polveri dell’AIM
• Introduzione alla metallurgia delle
polveri
• Le polveri metalliche: produzione e
proprietà
• Particolari sinterizzati: fasi della
produzione
• Indicazioni sulla progettazione delle forme
• Nuove tendenze della pressatura –
eco & energy
• Non destructive testing and computerized vision in PM parts production
• I materiali a base rame sinterizzati
• Metal Injection Molding e/o additive manufacturing
• Gli acciai inossidabili sinterizzati
• Il controllo della durezza e gli esami
microstrutturali sui sinterizzati
• Gli acciai speciali da polveri
• I metalli duri (carburi cementati)
• La pressatura isostatica a caldo
• Evoluzione delle polveri per la produzione di componenti sinterizzati.
Ottimizzazione delle miscele ed
elementi di lega
Per ulteriori informazioni rivolgersi
alla Segreteria dell’Associazione Italiana di Metallurgia, oppure visitare il
sito www.aimnet.it
Giornata di studio
LA METALLOGRAFIA
ATTRAVERSO UNA
CORRETTA PREPARATIVA
Milano, 26 maggio 2015
Le proprietà dei manufatti metallici
ed il loro comportamento in opera è
influenzato da molteplici fattori, chimici, fisici e meccanici.
La metallografia è la tecnica di analisi
della microstruttura dei metalli che
51
Atti e notizie
permette di valutarne le caratteristiche e il comportamento.
Mediante l’approfondimento metallografico dell’evoluzione delle
microstrutture - dal momento della
solidificazione delle leghe fino all’ottenimento del manufatto finale - è
possibile evidenziare i punti di forza
e/o le debolezze di determinati materiali, nonché comprendere come ottenere morfologie ottimali ed idonee
a garantire le migliori prestazioni dei
materiali in funzione dell’applicazione
a cui saranno assoggettati in opera.
Si può quindi affermare che competitività e qualità passino attraverso
l ottimizzazione delle microstrutture.
Molteplici campi di applicazione dei
metalli sarebbero impensabili senza
l’impiego di una corretta tecnica metallografica che va dall’identificazione
di fasi ed evoluzioni microstrutturali,
alla caratterizzazione di prodotti dopo
le lavorazioni, allo studio del danneggiamento meccanico o corrosionistico, e a mille altri aspetti.
Il Centro di Studio Trattamenti Termici e Metallografia dell’AIM, in collaborazione con STRUERS - azienda
leader nella produzione di strumentazioni e consumabili per le prepa52
razioni metallografiche - promuove
una Giornata di studio, che vedrà la
partecipazione straordinaria di George Vander Voort, eminente esperto a
livello mondiale sulla Metallografia ed
analisi correlate, principale autore di
oltre 370 articoli tecnici e di edizioni
riconosciute a livello mondiale.
Il programma della manifestazione
tratterà i seguenti argomenti:
• Come ridurre/ottimizzare i costi
di preparazione metallografica in
laboratorio
• Ispezione metallografica dopo trattamenti termochimici
• The Microstructure of Iron-Based
Alloys
• Fractography. Studying fracture
surface characteristics for better
failure diagnosis
• Studio sulla preparativa di un campione metallografico di qualità
• Etching experiences
sito www.aimnet.it.
Per tentare di dare un quadro dello stato dell’arte in questo campo il
Centro di Studio Acciaieria dell’AIM
ha pensato di organizzare una manifestazione specificatamente dedicata
a questo settore.
Dopo una presentazione sullo stato
dell’industria refrattaria italiana ed
europea presentata da Confindustria
Ceramica, è previstauna illustrazione
di carattere generale sulle failures dei
refrattari, sulle possibili cause e sui
metodi di analisi e successivamente
una memoria impiantistica sui principi di rivestimento del forno elettrico;
nel corso dei successivi interventi
verranno analizzate una per una problematiche e cause nei diversi reattori metallurgici per la produzione di
acciaio dalla fusione al colaggio.
In questo modo si pensa di poter
offrire a produttori ed utilizzatori un
proficuo scambio di discussione e
confronto per sempre meglio ottimizzare le scelte future.
Il programma dettagliato della Giornata – coordinata da M. Martino -
Giornata di Studio
FAILURES NEI
REFRATTARI
4 giugno 2015
Le problematiche relative alle failures
(cedimenti, malfunzionamenti, cattive
rese ecc.) dei prodotti refrattari utilizzati nel rivestimento degli impianti siderurgici da sempre costituiscono un
appassionante campo di discussione
tra produttori ed utilizzatori.
E’ evidente che questo tipo di problematiche possono avvenire non
solo per cause direttamente riferibili
alla qualità dei refrattari impiegati
ma a tutta una serie di possibili e
complesse interazioni tra il refrattario e le pratiche operative dell’acciaieria, le soluzioni di montaggio, la
conduzione dell’impianto, il tipo di
metallurgia adottata ecc.
Proprio per questo motivo si sta sempre più affermando la filosofia del
“best to fit” nella ricerca del miglior
compromesso possibile tra resa, costo e possibili inconvenienti.
prevede interventi su: •Andamento e statistiche del settore refrattario
•Failure analysis sui refrattari
•Influenza dei parametri di processo, sicurezza e manutenzione del
Vita associativa
tino EAF, criteri di progettazione e
principi del rivestimento refrattario
•Fermate ed incidenti nel forno elettrico: ma è solo responsabilità del
refrattario?
•“Failure” del lining in COV Limiti di
un lining non idoneo e analogie con
EAF
•Arresto siviera: cause e soluzioni
•Sollecitazioni sui rivestimenti di dolomite e soluzioni efficaci
•Criticità nei refrattari di placca per
il colaggio dell’acciaio
•Paniera: importanza della corretta scelta dei refrattari e della loro
messa in opera
sito www.aimnet.it
LE PROSSIME MANIFESTAZIONI AIM
SOLIDIFICAZIONE E COLATA CONTINUA
Corso itinerante – Centro A
7-8-14-15-22 maggio
Eur. Conf. HEAT TREATMENT & SURFACE ENGINEERING & 22nd
IFHTSE Congress
Venezia, 20-22 maggio
http:aimnet.it/ht2015.htm
LA METALLOGRAFIA ATTRAVERSO UNA CORRETTA
PREPARATIVA
Gds – Centro TTM
Milano, 26 maggio
TUBI IN ACCIAIO AL CARBONIO: Tubi saldati
GdS – Centro LP
Cremona, 25 giugno
FAILURES NEI REFRATTARI
GdS – Centro A
Milano, 4 giugno
SCUOLA DI METALLURGIA DELLE POLVERI
Quote sociali AIM 2015
(ANNO SOLARE)
Benemeriti (quota minima).1.750,00 €
Sostenitori (quota minima)... 750,00 €
Ordinari (solo persona)........... 70,00 €
Corso– Centro MP
Imola, 11-12 giugno
TUBI IN ACCIAIO AL CARBONIO: Tubi senza saldatura
GdS – Centro LP
Dalmine, 11 giugno
XI Giornate Nazionali CORROSIONE E PROTEZIONE
Ferrara, 15-16-17 giugno
http:aimnet.it/gncorr2015.htm
TUBI IN ACCIAIO AL CARBONIO: Tubi saldati
Seniores................................... 25,00 €
GdS – Centro LP
Cremona, 25 giugno
Juniores.................................... 15,00 €
METALLI A GRANO ULTRAFINE
La quota dà diritto di ricevere la rivista
dell’Associazione La Metallurgia Italiana. Ai soci viene riservato un prezzo
speciale per la partecipazione alle manifestazioni AIM e per l’acquisto delle
pubblicazioni edite da AIM.
Per ulteriori informazioni, iscrizioni,
rinnovi:
AIM
Associazione Italiana di Metallurgia
Piazzale R. Morandi, 2
20121 Milano
Tel.: 02 76021132/76397770,
fax: 02 76020551
Gds – Centri o MFM
Parma, 3 luglio
PROGETTAZIONE STAMPI
Corso Avanzato – Centro P
Bergamo, 9 – 10; 21 – 22 luglio
STEELSIM 2015
6th Int. Conf. Modelling and Simulation of Metallurgical Processes in
Steelmaking
Bardolino, 23-25 settembre
http:aimnet.it/steelsim2015.htm
TRATTAMENTI TERMICI
Corso modulare – Centro TTM
Milano, 29-30 settembre
Pillole per Preposti: LA MACCHINA FUSORIA
Corso – Centro A
Brescia, 14 ottobre
MATERIALI DI CARICA IN ACCIAIERIA
Gds – Centro A
Milano, 18 novembre
e-mail: [email protected]
www.aimnet.it
Per ulteriori informazioni rivolgersi alla Segreteria AIM, e-mail: [email protected],
oppure visitare il sito Internet www.aimnet.it
53
Atti e notizie
Attività dei Comitati Tecnici
CENTRO LAVORAZIONI PLASTICHE (LP )
(riunione del C.T. – 18 settembre 2014)
Manifestazioni in corso di organizzazione
- Ravanelli, Coordinatore della GdS “Fabbricazione dei tubi
senza saldatura”, ha comunicato che in linea di principio la
manifestazione può essere effettuata; si prevede di tenerla
a Dalmine a maggio, distribuita su due giornate, in funzione
dei contenuti, con visita agli stabilimenti.
- Donini e Fanchini, Coordinatori della Gds “Linee di
processo e finitura per prodotti piani” riferiscono della
difficoltà di trovare, in questo momento particolare della
siderurgia, una sede idonea a consentire anche la visita ad
impianti. Si conviene sulla opportunità di attendere che le
varie problematiche si risolvano, mantenendo monitorata
la situazione.
- Il Presidente Capoferri riferisce di quanto discusso in
comitato ristretto, con i Coordinatori Donini e.Gabrielli,
sugli esiti di una prima indagine conoscitiva in merito all’
ipotesi di promuovere una GdS o un Corso sull’ estrusione
dell’alluminio. In particolare si è convenuto di attendere
il convegno Alluminium 2000 – ICEB benchmark,
programmato a maggio a Firenze, dove confluiscono
tutti gli interessati al tema. Viene deciso di posticipare
l’iniziativa.
Iniziative future
- Si esamina l’eventualità di promuovere un’iniziativa del
tipo “Pillole per Preposti”, sperimentati con successo dal
Centro Acciaieria. Si individuano le criticità, tra le quali
quella di individuare gli impianti per le visite alle aziende
di interesse per il CT LP, che si prevede non abbiano
propensione ad ospitare. Concludendo, il Comitato rimanda
alle prossime riunioni la valutazione di come proseguire.
- Viene proposta una GdS dedicata al settore tiranteria,
dove il livello di conoscenze è molto basso. Lo stimolo a
partecipare può venire da interventi sulla nuova normativa
1090 sulle costruzioni, di difficile interpretazione e
applicazione senza adeguate conoscenze metallurgiche. Si
incarica Mariani di fare una prima indagine e di proporre
una possibile scaletta di contenuti alla prossima riunione.
- Si raccomanda di dedicare attenzione al settore
raccorderia oleodinamica, ove sono da poco in uso
particolari tipi di acciai con basso tenore di zolfo, che
presentano criticità di lavorazione. Viene chiesto a Mariani
di stilare una possibile scaletta di argomenti.
Stato dell’arte e notizie
- Si commenta che sulla rivista AIM compaiono articoli
scientifici, ma poco di carattere divulgativo. La cosa è
stata varie volte evidenziata, ma negli ultimi numeri la
tendenza sembra accentuata. Il Presidente ne parlerà in
sede di riunione dei Presidenti di Centro.
54
- Viene sollecitata anche una presa di posizione di AIM
per spiegare ai media cosa sia una acciaieria, in modo
da evitare preconcetti, che sono stati causa di accuse
infondate su personale tecnico innocente.
CENTRO METALLI E TECNOLOGIE APPLICATIVE
(MTA)
(riunione del C.T. – 9 dicembre 2014)
- Si informa che, in merito alla programmazione della
GdS. “L’impiego dell’acciaio nelle costruzioni civili”, è
stato deciso di posticipare l’appuntamento al 18 marzo
2015. Infatti i tempi tra l’annuncio e lo svolgimento della
Giornata risultavano troppo ristretti nel caso della prima
data individuata, col rischio di non riuscire a coinvolgere
tutti i possibili interessati. Soprattutto, non si sarebbe
potuta istruire la procedura per l’ottenimento dei crediti
formativi da parte dell’Ordine degli Ingegneri di Milano per
i partecipanti. La nuova data è stata accettata da tutti i
relatori, che hanno riconfermato la loro disponibilità.
Iniziative future
- Il Presidente Debernardi ha riassunto i principali temi
trattati nell’incontro dei Presidenti dei Centri di Studio AIM
con il Presidente AIM Mapelli, svoltasi il 27 novembre 2014.
In quella sede è emerso che per lo svolgimento delle attività
programmate nel 2015 è richiesta la collaborazione del CT
MTA per l’organizzazione di tre eventi. Con il Presidente del
Centro Fonderia, Caironi, sono già avvenuti contatti per la
GdS “ Fonderia delle leghe di rame”, organizzata da AIM
con Assofond. Da parte del Presidente del Centro Metalli
Leggeri, Vedani, è stata richiesta collaborazione, in fase di
studio, per una Giornata sul titanio. Il Presidente del Centro
Rivestimenti, Bestetti, propone poi di collaborare ad una
Giornata su ” Rivestimenti ed ingegneria delle superfici per
l’industria alimentare”.
CENTRO RIVESTIMENTI (R)
(riunione del C.T. – 15 dicembre 2014)
- Per quanto riguarda la preparazione del Corso
“Rivestimenti: modulo rivestimenti per via umida”, nel quale
verranno trattati anche i rivestimento sol-gel, si propone e
si approvano le date del 1 e 2 luglio 2015. Si conferma
la sede di Padova. I Coordinatori della manifestazione
saranno Bestetti, Bolelli, Brisotto. Viene approvato il
programma definitivo.
- Per la GdS “Rivestimenti ed ingegneria delle superfici
per l’industria alimentare”, si è accolta la disponibilità del
Presidente del Centro MTA a individuare presentazioni. Tra
i presenti si manifestano adesioni ad effettuare interventi.
Dai Centri
Parma si ipotizza come sede. Si discute poi una bozza della
locandina.
i docenti. Rimane da definire la tipologia specifica per la
presentazione dei casi pratici di failure.
Iniziative future
- Si decide di prevedere, per l’inizio del 2016, una Giornata
sulla caratterizzazione dei rivestimenti e delle superfici.
CENTRO TRATTAMENTI TERMICI E METALLOGRAFIA
Stato dell’arte e notizie
- Affrontando il rinnovo delle cariche, vengono eletti
all’unanimità come Presidente Bolelli; come Vicepresidente
Brisotto e come Segretario Bestetti.
- Si commenta il Convegno Nazionale sui trattamenti
termici e i rivestimenti, che si terrà a Venezia il 20-22
maggio. Si fa presente che l’80% delle memorie viene da
paesi stranieri.
- Si ricorda che dal 6 al 10 luglio UNIBS sta organizzando
una summer school sull’ALD. Per informazioni ci si può
rivolgere al Vicepresidente Brisotto.
CENTRO CONTROLLO E CARATTERIZZAZIONE
PRODOTTI (CCP)
(riunione del C.T. – 30 ottobre 2014)
- Il Corso “Prove Meccaniche” è stato programmato per
i giorni 24-26 febbraio 2015, presso il CSM di Roma
che fornisce i docenti, le aule, il materiale didattico e i
laboratori per le dimostrazioni, con un contributo spese
da parte di AIM. Il Coordinatore Trentini conferma che il
programma è ormai definito, con profondi cambiamenti
rispetto ai corsi precedenti; rimane confermata la durata
di tre giorni consecutivi, con l’ultima giornata riservata alle
visite e alle prove in laboratorio.
- Stella dà la sua disponibilità a coordinare il Corso
“Analisi chimiche” . Si concorda sulla necessità di rivedere
profondamente il programma classico, evitando per quanto
possibile l’intervento di produttori di strumentazione che
spesso forniscono presentazioni “commerciali”, con scarso
interesse didattico. Si curerà di evitare la sovrapposizione
con altre manifestazioni AIM che potrebbero interessare
gli stessi partecipanti. Verranno inseriti nuovi argomenti
che comprenderanno tecniche di analisi chimica nel
campo ambientale e della sicurezza e sui rivestimenti
superficiali.
Iniziative future
- L’organizzazione di una prevedibile GdS su circuiti
interlaboratorio è demandata ad ALPI in collaborazione con
RTM, che ha recentemente organizzato un Round Robin
sul Creep. La manifestazione è prevedibile per l’autunno
2015, con sede a Cormano o presso AIM, in funzione del
numero dei partecipanti.
- Tenuto conto che un Corso “Failure Analysis” è stato
appena proposto dal Politecnico di Milano con una
consistente partecipazione, si programma di organizzare
la 9° edizione del Corso AIM per l’autunno 2015. Il
coordinamento rimane a cura di Fossati. Il programma già
collaudato potrà presentare variazioni per quanto riguarda
(riunione del C.T. – 22 gennaio 2015)
- Il Presidente Petta illustra le decisioni riguardanti il Congresso “European Conference on Heat Treatment 2015 &
22°IFHTSE Congress” (20÷22 maggio 2015, Venezia Mestre). Essendo pervenute circa 170 memorie sono state
selezionate 90 di esse per le sessioni orali, cui si aggiungono due keynotes e un consistente numero di memorie
poster, applicando criteri di selezione basati sui contenuti
e sulla provenienza. La sessione poster avrà a disposizione
un ampio spazio e verrà conferito un premio al miglior poster. Al momento sono ipotizzati chairman singoli, ai quali
si chiederà di essere molto rigidi nel rispetto dei tempi di
esposizione per consentire ai partecipanti di seguire sessioni parallele.
- Il programma del Corso Modulare “Trattamenti Termici”
viene leggermente compattato per riuscire a contenere
l’attività in 4 moduli, ovvero otto giornate di lezioni più una
giornata di visita ad uno stabilimento (da accorpare all’ultimo modulo). Per individuare la sede di quest’ultimo sarà
da accertare la disponibilità degli ospitanti a dare l’accesso indistinto a tutti gli iscritti.
- Si discute relativamente al Corso “Metallurgia di Base” e,
valutato che l’iniziativa comporta un impegno notevole, si
ipotizza di posticiparlo al 2016.
- Il Coordinatore della GdS “Trattamento termico dei sinterizzati”, Morgano, ha fatto pervenire una bozza di programma
che prevede la collaborazione del Centro Metallurgia delle
Polveri e l’effettuazione a fine settembre/inizio ottobre, in
sede da definire. I contenuti saranno incentrati sulle applicazioni automotive e si terrà una tavola rotonda conclusiva,
eventualmente con inviti mirati a specialisti del settore.
- Bavaro, Coordinatore della GdS “Contributo della metallografia alla Failure Analysis” sta contattando i relatori della precedente edizione e raccogliendo nuove adesioni per
inserire argomenti diversi. Si concorda di puntare a collocare la manifestazione a novembre. Durante la prossima
riunione verrà discussa una prima bozza di programma.
Iniziative future
- Per il 2016, oltre a rispettare le cadenze tradizionali dei
Corsi, si ipotizzano le seguenti attività: GdS su trattamenti
criogenici / tecniche di tempra innovative; GdS su materiali per lo stampaggio (incentrato sull’automotive); GdS
“Trattamenti superficiali senza materiale d’apporto” (ad
esempio: tempra a induzione, laser, plasma, ecc). Si fa una
prima valutazione anche di una GdS dedicata a tecniche di
tempra innovative, nella quale inserire l’argomento “trattamenti criogenici”, sui quali sono in corso diversi lavori
scientifici e che si stanno diffondendo all’interno di molte
aziende.
55
Normative
Atti e notizie
Notizie da UNSIDER
Norme pubblicate e progetti in inchiesta (aggiornamento 7 aprile 2015)
Norme sui metalli non ferrosi
pubblicate da UNI nel marzo 2015
Progetti al voto FprEN e ISO/FDIS
- aprile 2015
UNI EN 1559-4:2015
FprEN – progetti di norma europei
Fonderia - Condizioni tecniche
di fornitura - Parte 4: Requisiti
addizionali per getti di leghe di
alluminio
FprEN ISO 13702
Progetti in inchiesta prEN e ISO/
DIS - aprile 2015
prEN – progetti di norma europei
prEN ISO 16530-1
Petroleum and natural gas industries.
Well integrity. Part 1: Life cycle
governance (ISO/DIS 16530-1:2015)
Petroleum and natural gas
industries. Control and mitigation
of fires and explosions on
offshore production installations.
Requirements and guidelines (ISO/
DIS 13702:2013)
FprEN ISO 24817
Petroleum, petrochemical and
natural gas industries. Composite
repairs for pipework. Qualification
and design, installation, testing and
inspection (ISO/FDIS 24817:2015)
prEN ISO 19901-4
Specific requirements for offshore
structures. Part 4: Geotechnical and
foundation design considerations
(ISO/DIS 19901-4:2015)
FprEN ISO 16961
natural gas industries. Internal
coating and lining of steel storage
tanks (ISO/FDIS 16961:2015)
prEN ISO 17781
FprEN ISO 19901-8
natural gas industries. Test methods
for quality control of microstructure
of austenitic/ferritic (duplex)
stainless steel (ISO/DIS 17781:2015)
Petroleum and natural gas
industries. Specific requirements for
offshore structures. Part 8: Marine
soil investigations (ISO 199018:2014)
prEN 16808
FprCEN ISO/TS 17969
natural gas industries. Safety of
machineries. Manual elevators
natural gas industries. Guidelines on
competency for personnel (ISO/DTS
17969:2015)
ISO/DIS – progetti di norma
internazionali
ISO/DIS 16530-1
Well integrity. Part 1: Life cycle
governance
56
FprEN 1754
Magnesium and magnesium alloys.
Designation system for anodes,
ingots and castings. Material
symbols and material numbers
ISO/FDIS – progetti di norma
internazionali
ISO/FDIS 4689-2
Iron ores. Determination of sulfur
content. Part 2: Combustion/
titration method
ISO/FDIS 4689-3
Iron ores. Determination of sulfur
content. Part 3: Combustion/
infrared method
ISO/FDIS 10203
Iron ores. Determination of
calcium. Flame atomic absorption
spectrometric method
ISO/FDIS 10204
magnesium. Flame atomic
absorption spectrometric method
ISO/FDIS 11536
Iron ores. Determination of loss on
ignition. Gravimetric method
ISO/FDIS 15633
nickel. Flame atomic absorption
spectrometric method
ISO/FDIS 15634
chromium content. Flame atomic
absorption spectrometric method
ISO/FDIS 24817
natural gas industries. Composite
repairs for pipework. Qualification
and design, installation, testing and
inspection
Giornate Nazionali
sulla
Corrosione
e Protezione
XI EDIZIONE
Ferrara - 15-17 giugno 2015
www.aimnet.it/gncorr2015.htm
Organizzate da
ASSOCIAZIONE
ITALIANA
DI METALLURGIA
Con il patrocinio di
Provincia di Ferrara
Coordinatore delle Giornate
Prof. Cecilia Monticelli
Presentazione
Le Giornate Nazionali sulla Corrosione e Protezione si terranno a Ferrara dal 15 al 17 giugno 2015. Si tratta di un
evento che negli anni si è saputo affermare su scala nazionale come punto di incontro per discutere questioni
scientifiche, tecnologiche e produttive, nell’ambito della corrosione e protezione dei materiali. Il Convegno
prevede la presentazione dei risultati raggiunti da vari gruppi di studio e da numerose aziende del settore.
Anche in questa undicesima edizione sono stati istituiti dei premi, destinati a giovani ricercatori che si distingueranno, nell’ambito della manifestazione, per l’importanza e l’attualità dei temi proposti nelle loro letture.
Aree tematiche principali
• Corrosione negli ambienti naturali: acque, atmosfera, terreno
• Corrosione negli impianti e nelle strutture industriali
• Tecniche di studio e controllo dei fenomeni corrosivi
• Corrosione e protezione delle armature nelle opere in c.a.
• Inibitori di corrosione
• Case histories
• Corrosione nei beni culturali
• Corrosione in ambiente biologico
• Protezione catodica
• Rivestimenti e trattamenti superficiali
Spazio aziende
È previsto uno spazio per l’esposizione di apparecchiature, per la presentazione dei servizi e per la distribuzione di materiale promozionale. Informazioni più dettagliate possono essere richieste alla Segreteria organizzativa del convegno
([email protected]).
Presentazione di memorie
Gli interessati a presentare memorie scientifiche dovranno inviare entro il 27 febbraio 2015, il titolo della memoria,
i nomi degli autori della memoria e la loro affiliazione ed un sommario di circa 500 parole mediante il modulo online
presente sul sito www.aimnet.it/gncorr2015.htm.
Date importanti:
Invio titolo e riassunti
Notifica accettazione
Apertura iscrizioni
Invio dei testi completi
27 febbraio 2015
20 marzo 2015
27 marzo 2015
30 aprile 2015
Atti
Gli atti del Convegno saranno predisposti sotto forma di CD-Rom e distribuiti agli iscritti all’inizio dei lavori.
Sede
La manifestazione si terrà dal 15 al 17 giugno presso le sale Imbarcadero del Castello Estense di Ferrara in Largo Castello
1 - Ferrara. Per maggiori informazioni consultare il sito: www.castelloestense.it.
Segreteria organizzativa
AIM - Associazione Italiana di Metallurgia
Piazzale Rodolfo Morandi 2 · 20121 Milano · Tel. 0276021132 / 0276397770 · Fax 0276020551
E-mail: [email protected] · Website: www.aimnet.it/gncorr2015.htm
Plug and Work –
Automazione e Simulazione
Visit us at
METEC 2015
Hall 5, Booth E22
GIFA/THERMPROCESS 2015
Hall 10, Booth H41
June 16 - 20, Düsseldorf, Germany
Provato e testato molto tempo prima dell’installazione
„Plug and Work“, è il prodotto più consolidato dai
nostri specialisti di Automazione. Dimenticate le lunghe
fasi di prova sui vostri impianti – il nostro sistema è collegato
ad un modello di simulazione che riproduce l’intero processo.
E’ così che noi proviamo accuratamente ed ottimizziamo
i programmi di lavoro e le funzioni tecnologiche prima della
messa in servizio.
Il risultato: sistemi di automazione testati e perfettamente
affidabili, che funzionano senza problemi. Benefici: curve
di avviamento ripide per i nuovi impianti e ammodernamenti.
Tempi brevi di assiemaggio che ottimizzano il ritorno sugli
investimenti. I vostri clienti potranno così beneficiare di una
produzione affidabile e di un rispetto dei termini di consegna.
SMS INNSE S.p.A.
Via Milano, 4
20097 San Donato Milanese (MI), Italy
Phone: +39 02 2124-1
Fax:
+39 02 2124-699
E-mail: [email protected]
Internet: www.innse.com

La Metallurgia Italiana

Transcription

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